Ageing influence on nickel-based superalloys at ...1022216/FULLTEXT01.pdf · these superalloys have...

MASTER’S THESIS

2006:251 CIV

ÅSA MARTINSSON

Ageing Infl uence on Nickel-basedSuperalloys at IntermediateTemperatures (400–600°C)

MASTER OF SCIENCE PROGRAMMEEngineering Materials • EEIGM

Luleå University of TechnologyDepartment of Applied Physics and Mechanical Engineering

Division of Engineering Materials

2006:251 CIV • ISSN: 1402 - 1617 • ISRN: LTU - EX - - 06/251 - - SE

ABSTRACT

Nickel-based superalloys are employed in a variety of gas turbine components due to their excellent high-temperature properties. Unfortunately, several of these superalloys have proved to shrink during long-term service/ageing at intermediate temperatures (400-600°C). This study comprises an investigation of the ageing effects of two nickel-based superalloys from different classes of material; Haynes 230 and single crystal CMSX-4. The work is divided into three parts:

I. Literature study. Haynes 230 and CMSX-4 were studied as well as reports and articles concerning shrinkage and microstructural changes owing to ageing in different nickel-based superalloys.

II. Mechanical testing. The ageing effects on mechanical and physical properties were investigated in a series of experiments. The employed methods are tensile testing, impact testing, low cycle fatigue testing and hardness testing.

III. Microstructural examination. The underlying causes to the property changes were investigated in a microstructural examination including carbide size measurements, chemical analysis of the composition and analysis of micro images taken with scanning electron microscope and light optical microscope.

The results from the experimental testing show that the ageing has a marked effect on the properties of both CMSX-4 and Haynes 230. The low cycle fatigue properties of Haynes 230 are unaffected by the ageing while the yield and tensile strength and hardness are improved. The ductility is improved in both Haynes 230 and CMSX-4. The hardness of CMSX-4 was unaffected by the ageing while the tensile strength was slightly impaired. It was also found that the properties and the ageing effect depended on the single crystal direction. The study shows that the ageing has a significant effect on the precipitation of secondary carbides at grain boundaries in Haynes 230, which has a strengthening effect on the material. No distinct changes of the microstructure explaining the property changes in CMSX-4 were found. It is reported in other studies that shrinkage in other superalloys is caused by short and long range ordering based on Ni2Cr and Ni3Cr. The microstructural examination neither confirms nor refutes that the shrinkage and property changes are caused by this type of ordering. However, the results from the examination indicate that ordering formation is possible.

SAMMANFATTNING

Nickelbaserade superlegeringar används i ett antal gasturbinkomponenter på grund av deras utmärkta högtemperaturegenskaper. Tyvärr har ett flertal av dessa superlegeringar visat sig krympa vid långtidsåldring/användning vid medelhöga temperaturer (400-600°C). Denna studie omfattar en undersökning av åldringseffekterna för två nickelbaserade superlegeringar från olika materialklasser; Haynes 230 och CMSX-4. Arbetet är indelat i tre delar: IV. Litteraturstudie. Haynes 230 och CMSX-4 studerades liksom rapporter och

artiklar om krympning och förändringar i mikrostrukturen orsakade av åldring i olika nickelbaserade superlegeringar.

V. Mekanisk provning. Åldringseffekter på mekaniska och fysikaliska egenskaper undersöktes i en serie experiment. De använda metoderna är dragprovning, slagprovning, lågcykelutmattning och hårdhetsprovning.

VI. Mikrostrukturundersökning. De bakomliggande faktorerna till förändring av egenskaperna undersöktes i en mikrostrukturstudie som omfattade mätning av karbidstorlek, kemisk analys av sammansättningen och analys av foton tagna med svepelektronmikroskop och optiskt ljusmikroskop.

Resultaten från den experimentella delen visar att åldringen har markant effekt på egenskaperna för både CMSX-4 och Haynes 230. Lågcykelutmattningsegenskaperna är oförändrade efter åldring i Haynes 230 medan sträck- och brottgräns samt hårdheten förbättras. Duktiliteten förbättras i både Haynes 230 och CMSX-4. Hårdheten i CMSX-4 påverkas inte av åldringen medan dragbrottstyrkan försämrades något. Resultaten visar också att egenskaperna och åldringens påverkan beror på enkristallens riktning. Studien visar att åldringen har en påfallande effekt på utskiljningen av sekundära karbider i korngränserna i Haynes 230, vilket har en stärkande effekt på materialet. Inga påtagliga ändringar i mikrostrukturen som förklarar förändringen av egenskaperna i CMSX-4 hittades. Det har rapporterats i andra studier att krympning i andra superlegeringar orsakas av ”short and long range ordering” baserad på Ni2Cr and Ni3Cr. Mikrostrukturstudien varken bekräftar eller motbevisar att krympningen och förändringen av egenskaperna beror på den här typen av ordning i materialen. Resultaten indikerar dock att den här typen av ordning är möjlig.

PREFACE

The work presented in this Master’s thesis was carried out at Siemens Industrial Turbomachinery AB, Finspång, during the period Mars –August 2006. To the persons who made this thesis possible, I give my deepest thanks. I would like to express my sincere gratitude to my supervisor Johan Moverare (Division of Materials Technology, SIT AB) for his great support and all the fruitful discussions we have had, for his encouragement and for believing in me. A special thank goes to Carina Jansson who was a great help in the laboratory and Ramón Niño Lopez for having an answer to all my questions. I would also like to thank everybody at the Division of Materials Technology for making me feel like home during my time at SIT AB. Finally I thank Magnus Odén (Division of Engineering Materials, Luleå University of Technology) for his time and guidance.

NOMENCLATURE

A5 Elongation of waist of tensile test specimen at% Atom percent Bal Balance BEI Backscattered electron image D Diameter of waist of tensile test specimen after test D0 Initial diameter of waist of tensile test specimen d1 Mean diameter of primary carbides d2 Mean diameter of secondary carbides DS Directionally solidified E Elastic modulus EDS Electron dispersive spectroscopy EDX Electron dispersive X-ray FCC Face centred cubic HCF High cycle fatigue K´ Strength coefficient KV Energy absorbed during impact test L Length of waist of tensile test specimen after test L0 Initial length of waist of tensile test specimen LCF Low cycle fatigue LOM Light optical microscope LRO Long range ordering N Number of cycles in LCF test n´ Strain-hardening exponent Nf Number of cycles to failure in LCF test, 50 % load drop or failure Ni Number of cycles to crack initiation in LCF test, 5 % load drop r1 Aspect ratio of primary carbides r2 Aspect ratio of secondary carbides ref Reference Rm Tensile strength Rp0.1 Yield strength, calculated at 0.1 % of plastic strain deformation Rp0.2 Yield strength, calculated at 0.2 % of plastic strain deformation RT Room temperature SEI Secondary electron image SEM Scanning electron microscope SIT AB Siemens Industrial Turbomachinery AB SRO Short range ordering SX Single crystal TBC Thermal barrier coating TC Critical temperature TCP Topologically close-packed temp Temperature vol% Volume percent wt% Weight percent Z Reduction of cross-section area of tensile test specimen γ Gamma matrix phase γ´ Gamma prime phase Δεpl Plastic strain range ρ Electrical resistivity σ phase A specific type of topologically close-packed phase σmax Maximum stress

CONTENTS

1 INTRODUCTION 9

1.1 SIEMENS INDUSTRIAL TURBOMACHINERY AB 9 1.2 BACKGROUND 9 1.3 PURPOSE 9 1.4 PRESENTATION OF THE PROBLEM 9 1.4.1 Objective 10 1.5 DELIMITATIONS 11 1.6 THEORETICAL BACKGROUND 11 1.6.1 Gas turbines 11

1.6.1.1 The compressor 12 1.6.1.2 The combustion chamber 12 1.6.1.3 The turbine 12

1.6.2 Turbine blade 13 1.6.3 Combustion chamber 14 1.6.4 Equipment 15

1.6.4.1 Tensile test equipment 15 1.6.4.2 Impact test equipment 16 1.6.4.3 Low cycle fatigue test equipment 16 1.6.4.4 Scanning electron microscope (SEM) 17 1.6.4.5 Light optical microscope 18 1.6.4.6 Stereo microscope 19 1.6.4.7 Hardness test equipment 19

2 SUPERALLOYS 20

2.1 NICKEL-BASED SUPERALLOYS 20 2.1.1 Composition 20 2.1.2 Phases and microstructure 20

2.1.2.1 Gamma matrix (γ) phase 21 2.1.2.2 Gamma prime (γ´) phase 21 2.1.2.3 Carbides 22 2.1.2.4 Grain boundary (γ´) phase 22 2.1.2.5 Borides 23 2.1.2.6 Topologically close-packed (TCP)-type phases 23

2.1.3 Strengthening of superalloys 23 2.1.3.1 Solid-solution strengthening 23

2.1.3.2 Precipitation hardening 24 2.1.4 Rafting 24 2.2 GAS TURBINE BLADE CASTING 25 2.2.1 Directional solidification (DS) 25 2.2.2 Single crystal (SX) casting 26 2.3 SINGLE CRYSTAL SUPERALLOY CMSX-4 26 2.4 SUPERALLOY HAYNES 230 27

3 AGEING EFFECTS IN SUPERALLOYS 29

3.1 CONTRACTION IN SUPERALLOYS 29 3.2 SHORT RANGE ORDERING (SRO) AND LONG RANGE ORDERING (LRO) 32 3.3 NEGATIVE CREEP 35 3.4 CARBIDE PRECIPITATION 35

4 AGEING INFLUENCE ON MECHANICAL PROPERTIES AND MICROSTRUCTURE OF HAYNES 230 AT INTERMEDIATE TEMPERATURES (500-550°C) 37

4.1 SPECIMEN PREPARATION AND EXPERIMENTAL PROCEDURE 37 4.1.1 LCF test 37 4.1.2 Impact test 37 4.1.3 Tensile test 37 4.1.4 Microscopic examination 38 4.1.5 Hardness test 39 4.2 TEST RESULTS 40 4.2.1 LCF test 40 4.2.2 Impact test 42 4.2.3 Tensile test 43 4.2.4 Hardness test 45 4.2.5 Carbide size measurements 47 4.2.6 Chemical analysis 48 4.3 DISCUSSION 50 4.3.1 Mechanical tests 50 4.3.2 Microstructural examination 52 4.4 CONCLUSION 55

5 AGEING INFLUENCE ON MECHANICAL PROPERTIES AND MICROSTRUCTURE OF CMSX-4 AT INTERMEDIATE TEMPERATURES (500-550°C) 56

5.1 SPECIMEN PREPARATION AND EXPERIMENTAL PROCEDURE 56

5.1.1 Tensile test 56 5.1.2 Microscopic examination 56 5.1.3 Hardness test 57 5.2 MECHANICAL TEST RESULTS 58 5.2.1 Tensile test 58 5.2.2 Angle of slip plane 61 5.2.3 Hardness test 61 5.3 DISCUSSION 62 5.3.1 Mechanical tests 62 5.3.2 Microstructural Examination 63 5.4 CONCLUSION 65

6 FUTURE WORK 67

7 REFERENCES 68

8 ANNEXES 70

8.1 MATERIAL DATASHEET FOR HAYNES 230 COATING 70 8.2 TEST CERTIFICATE, HAYNES 230, HEAT NO 1830587801 71 8.3 DRAWINGS 72 8.3.1 Low cycle fatigue test specimen 72 8.3.2 Impact test specimen 73 8.3.3 Tensile test specimen 74 8.4 CALCULATIONS 75 8.4.1 LCF test 75 8.4.2 Tensile test 75 8.5 RESULTS 76 8.5.1 Tensile test, Haynes 230 76 8.6 MICROSCOPY PHOTOS 77 8.6.1 Haynes 230 77 8.6.2 CMSX-4 89

9

1 INTRODUCTION

1.1 SIEMENS INDUSTRIAL TURBOMACHINERY AB

Siemens is one of the world’s largest electrical engineering and electronics companies. Some of the areas where the company is active are Information and Communications, Automation and Control, Power, and Medical. Siemens is represented in 190 countries all over the world. Siemens Industrial Turbomachinery AB (SIT AB) is a part of the business sphere Power Generation and is by far the largest of Siemens companies in Sweden with a turnover of 6 billion SEK. SIT AB is located in Finspång and Trollhättan and has over 2000 employees working to develop, produce and install gas turbines, steam turbines and power plants for industrial application. The effect of the gas turbines is in the area 15-50 MW and the steam turbines are in the area 60-180 MW.

1.2 BACKGROUND

The largest type of gas turbines produced in Finspång is SGT-800 with an electrical output of 45 MW and an electrical efficiency of 37 %. The large electrical output requires very high temperatures in operation in the combustion chamber as well as in the turbine. The high-temperature properties of steel are not good enough for these components, so different types of nickel-based superalloys are used. One of these superalloys is Haynes 230, used in the combustion chamber. The sheets surrounding the flames can reach a temperature of over 900°C as a maximum. The outer parts of the combustion chamber, previously made of Hastelloy X, which belongs to the same class of material as Haynes 230, reach a temperature not higher than approximately 500°C. CMSX-4 is a nickel-based superalloy of a different class, used in the hottest part of the turbine. The airfoil of the blade in the first turbine stage is exposed to hot gases, heating the material to a temperature of up to 900°C. On the other hand, the temperature in the blade root is within the temperature range 400-600°C.

1.3 PURPOSE

The purpose of this diploma work is to examine microstructural changes in nickel-based superalloys due to ageing at intermediate temperatures (400-600°C). The study will also comprise an examination of mechanical and physical properties in order to find any possible changes owing to ageing and phase transformation.

1.4 PRESENTATION OF THE PROBLEM

The use of nickel-based superalloys isn’t uncomplicated. Despite their excellent high-temperature strength, several superalloys have proved to shrink at intermediate temperatures, 400-600°C [1]. The results from these measurements

10

will be discussed in detail in chapter 3. Haynes 230 and Hastelloy X are two examples. The affected part is the combustor mounting ring which makes it possible to bolt the combustor chamber to the turbine casing. This ring with bolt holes used to be made of Hastelloy X, but this material caused problems during inspection. The ring has to be removed during inspection of the combustion chamber and afterwards it didn’t fit because of shrinkage. A short-term solution of the problem has been to make the ring out of Cr-Mo steel instead. The shrinkage measurements carried out at Siemens [1] proved that Haynes 230 was the superalloy in the study that exhibited the largest shrinkage and at the shortest time during ageing. This is why it was chosen to be examined experimentally in this study. It is desired to understand the microstructural changes that are behind the shrinkage phenomenon. Previous studies of simple binary Ni-Cr systems and tertiary Ni-Cr-Fe systems prove that ageing of these materials can result in a long range ordering of the structure [2]. Whether this ordering phenomenon exists in more complex alloys such as commercial nickel-based superalloys isn’t established but there are indications that this could be the case. A phase transformation is likely to affect the mechanical properties of the alloy [2]. This could for example affect the fracture behaviour and thereby shorten the life time of the components. No measurements in order to determine the existence of shrinkage in CMSX-4 have been carried out so far, but the shrinkage behaviour has been observed in similar alloys at intermediate temperatures [3]. Since the service temperature of the blade root is within the critical temperature range, the material was included in this study. If the shrinkage phenomenon affects the superalloy, hopefully the results of this study will facilitate the life time assessment and the dimensioning of the turbine blade.

1.4.1 Objective

The aim of this diploma work is to map out and analyse any changes in structure and mechanical properties between new and aged material in order to gain a deeper understanding of the behaviour of Ni-based superalloys. The experimental work that ought to be done is:

1. to perform tensile test, impact test, low cycle fatigue test and hardness test on Haynes 230

2. to perform tensile test and hardness test on CMSX-4 3. to examine the fracture surface of the tensile test specimens 4. to examine the cross-section of the tensile test specimens (scanning

electron microscope and light optical microscope) 5. to analyse the composition of the matrix and the carbides in Haynes 230 6. to measure the carbide size in Haynes 230 7. to analyse existing experimental data about negative creep from a

metallurgical perspective The first six paragraphs are prioritised and the last one will be carried out as far as time admits.

11

Figure 1: SGT-800 gas turbine [4].

1.5 DELIMITATIONS

Haynes 230 is examined in this diploma work because of its ageing behaviour. It’s the austenitic material found in a gas turbine that exhibits the largest and fastest shrinkage when heat treated at intermediate temperatures. Hastelloy X that is mentioned several times in the report exhibits a pronounced shrinkage as well, but the lack of material prevented the incorporation of the alloy in the experimental part of this study. However, Hastelloy X is very similar to Haynes 230 in composition and behaviour, and it’s therefore reasonable to believe that the causes that lie behind the shrinkage phenomenon are the same for both materials.

1.6 THEORETICAL BACKGROUND

1.6.1 Gas turbines

Gas turbines are power generating engines used for a variety of industrial applications, such as aircraft and marine propulsion, driving generators or pumps and electricity generation in power plants. The three main components in a gas turbine are the compressor, the combustion chamber and the turbine. A cross-section diagram of gas turbine SGT-800 is shown in Figure 1.

The gas turbine cycle is usually described with help of the ideal Brayton cycle, illustrated in Figure 2. Fresh air at atmospheric pressure flows into the compressor. The compression is isentropic and results in a rise in temperature and pressure. The compressed air is mixed with fuel and ignited in the combustion chamber, causing a dramatic temperature rise. The hot gases are then passing through the turbine where work is derived during the expansion of the gases. [5], [6]

12

Figure 2: Simplified scheme of the gas turbine system and the ideal Brayton cycle [7].

1.6.1.1 The compressor

Fresh air, holding ambient temperature and atmospheric pressure, enter the gas turbine under free stream conditions at station 1. Ideally, an isentropic compression takes place between station 1 and 2. The compression results in a volume reduction and a significant increase in pressure and temperature. In reality, the entropy isn’t unaffected by the compression. The work executed on the gas will slightly increase the entropy, making the line between station 1 and 2 in the T-s diagram in Figure 2 incline somewhat to the right. [6] In SGT-800 the air is compressed in 15 stages before the flow exits the compressor. The gas pressure is 19.3 atm after all stages of compression and the temperature of the gas exiting the compressor is approximately 430°C. [4], [8] The aerodynamics of the compressor is of highest importance for the efficiency of the gas turbine. Continuous efforts are made to improve the design of the compressor with the intention of minimising the consumption of work required during the compression of the air.

1.6.1.2 The combustion chamber

The hot air is mixed with fuel in the annular combustion chamber before ignition. The pressure is remained constant during the combustion (between station 2 and 3), but the heat added in the process will cause a critical increase in temperature. Because of the ideal gas law,

nRTPV = , Equation 1: Ideal gas law.

the volume will increase as well. The temperature of the hot exhaust depends on the type of fuel mixed with the compressed air but also on the proportion of fuel in the gas mixture [5], [6]. The fuel used in general is diesel oil or natural gas. The temperature of the exhaust after combustion is approximately 1420°C in SGT-800 [8].

1.6.1.3 The turbine

In the last process, the hot gases are allowed to expand isentropically while passing through the turbine. Between station 3 and 4, the temperature and the

13

pressure decreases over several turbine stages while work is derived from the expansion. The compressor and the turbine are mounted on the same shaft and a part of the work obtained in the turbine is used to drive the compressor. The exceeding work is either obtained in the form of mechanical drive used for driving external load, such as pumps or propellers, or in the form of electrical power. [5], [6] The gas enters the turbine, passing by the first stage vanes. The vanes are directing the gas in the optimal angle towards the following blades. The work performed by the gas is forcing the blade disk to rotate, from where the electrical power is derived. The gas continues towards the next set of vanes and the procedure is repeated. For each stage, the volume is increasing and the temperature is decreasing. The turbine outlet temperature is 546°C [4]. The hot gases can be used for heating water in a combined cycle with a steam turbine, which will diminish the energy losses and increase the efficiency notably.

1.6.2 Turbine blade

The most complex components in a gas turbine are found in the hot gas path. The turbine blades and vanes belong to these components. The turbine blade consists of two parts: the airfoil which is exposed to the flow of heat gases, and the root section which attaches the blade to the turbine disk. A picture of the first stage turbine blade from SGT-800 is found in Figure 3. The blades are exposed to very high temperatures and centrifugal stresses during service, facts that make great demands on the material and the design of the blade. The high temperature and the flow of contaminated exhaust expose the blades to oxidation and corrosion [9]. The high rotational speed of the turbine creates another problem. High centrifugal forces will affect the blade and lead to high stress in the material, which in turn result in creep. Other aggravating circumstances for the lifetime of the blade are the start-ups and stops. Great changes in temperature and unequal distributed loads will result in thermal strain in the material. [11] The design of the blade is worked out from the thermal and mechanical conditions in the turbine. We can take the case of the first stage blade in SGT-800 to illustrate the design of a blade. The gas temperature after the first vane is approximately 1180°C. The material in the first stage blade is CMSX-4, which has a melting temperature range between 1320°C and 1380°C [12]. In order to avoid high metal temperature, two measures are taken. The blade is cooled

Figure 3: First stage blade in SGT-800 [10].

Platform

Trailing edge

Airfoil Leadingedge

Root section

14

internally with air. The air is passing via the root section and out through the cooling holes. A thin film of cool air is then created on the outside of the blade, protecting it from the hot gases. Furthermore, the airfoil and the platform of the blade are covered with thermal barrier coating, TBC, in order to keep the metal temperature down. In this specific case the blade is coated with platinum aluminide to improve the oxidation and corrosion resistance. It’s important that the cooling doesn’t provoke a thermal gradient in the material since it would have a negative influence on the performance of the blade. The average temperature of the surface of the first stage blade is 850-900°C. The cooling is concentrated to the load-bearing parts where the temperature can be as low as <750°C. The blade tip belongs to the hotter areas where the local temperature can reach over 1000°C. [8] The root section holds a temperature of around 500°C. Because of the severe conditions, the lifetime of a turbine blade is today 20 000 operating hours. Continuous development of the material and the design of the blades and the vanes is performed in order to prolong the lifetime of the components and in that way reduce the cost.

1.6.3 Combustion chamber

The combustor section consists of three main parts: the fuel injection system, the combustion liner and the combustion outlet. See illustration of SGT-800 combustion section in Figure 4. The combustion chamber is the first part in the hot gas path. The components have to face loads at high temperatures when the mixture of compressed air and fuel are burned. The conditions in the combustion chamber are severe. The annular combustion chamber in SGT-800 is built up by an inner and outer wall, also called combustion liner. The combustion liner is constructed by sheets of Haynes 230, surrounding the flames. The walls are exposed to creep and oxidation due to high temperature and pressure in the chamber. Another influence is created by the flames’ tendency to pulsate. The intensity variation of the flames creates high cycle fatigue, HCF, vibrations in the walls. To sustain the strain, it takes a ductile material with maintained good strength and resistance to corrosion at high temperatures.

15

SGT-800 is the gas turbine with the highest firing temperature manufactured at SIT AB. The outlet gas temperature is, as mentioned before, approximately 1420°C [8]. Since the melting temperature range of Haynes 230 is 1301-1371°C [13] measures have been taken to keep the metal temperature at a considerably lower level. The sheets are coated with TBC, in this case in two layers. The bond coat is made up of NiCoCrAlY and the top coat consists of Y2O3-stabilised ZrO2. See Annex 8.1 for further information. Furthermore, film cooling is used to additionally lower the metal temperature. These steps result in an average metal temperature of the combustion liner between 800°C and 850°C. Near the flame and at the outlet of the combustion chamber, local temperatures above 900°C have been measured.

1.6.4 Equipment

1.6.4.1 Tensile test equipment

The tensile tests were performed in a Schenck-Trebel RSA 100 test machine. The test rig is controlled digitally through an EDC 120 controller. The test equipment consists of a clamping system and a load cell attached to a crosshead. The test rig has a loading capacity of 100 kN in both traction and compression tests. A photo of the equipment is found in Figure 5. An extensometer of the type Sandner with a gauge length of 10 mm was attached to the specimens to measure the strain during the first part of the test.

Figure 5: Schenck-Trebel RSA 100 tensile test equipment.

Figure 4: Combustor section, SGT-800.[4].

16

The test machine is also equipped with a furnace with a heating capacity of up to 1000°C and a measuring instrument for acoustic emissions. These tools weren’t used in this study.

1.6.4.2 Impact test equipment

The impact strength of Haynes 230 was investigated by using a Charpy impact test machine manufactured by Wolpert. The maximum energy possible to absorb is 150 J or 300 J, depending on the applied load. A photo of the equipment is found in Figure 6.

1.6.4.3 Low cycle fatigue test equipment

The universal testing machine Instron 1275 was used to carry out the strain controlled low cycle fatigue (LCF) tests. It’s a servo-hydraulic testing machine which can be used for dynamic and static tensile testing as well as for low cycle fatigue and high cycle fatigue testing. The test rig has a loading capacity of ± 100 kN. Equipped with the furnace visible in Figure 7, testing at temperatures up to 1000°C can take place. The test machine is controlled by the Instron fast track 8800 equipment. A dynamic extensometer from Instron with catalogue number 2620-602 was used to measure the strain. The extensometer has a gauge length of 12.5 mm and a travel of ± 2.5 mm. All the experiments in this study were effectuated at room temperature, RT, and in total strain control. Figure 7: Instron 1275 LCF equipment.

Figure 6: Charpy impact test equipment.

17

1.6.4.4 Scanning electron microscope (SEM)

The scanning electron microscope (SEM) is today a common tool used in a variety of branches in the industry. It’s used to reproduce surfaces and objects at high magnitude and to determine the composition of the material, both quantitatively and qualitatively. [14] A SEM consists of three main components: an electron gun, lenses and an electron detector. A schematic image of SEM is shown in Figure 8. The electron gun contains a filament, usually made of tungsten. When the filament is heated with an electrical current, a thermal emission of electrons is created. The electron gun accelerates the electron beam to an energy between 2 keV and 40 keV. Magnetic lenses are focusing the beam to a diameter of 1-10 nm before it reaches the sample. [14], [16] The scanning coils are moving the electron beam in a rectangular set of straight lines, also known as raster, over the sample. When the beam hits the sample, an interaction occurs between the surface of the sample and the incident electrons, which results in an emission of secondary electrons from the surface. A detector counts the number of secondary electrons. By comparing the intensity of the secondary electrons and the intensity of the incident electron beam, a topographic secondary electron image (SEI) is created. The magnitude of the image is easily changed by changing the size of the raster. [14], [16] When bombarded with electrons, the surface emits not only secondary electrons but backscattered electrons and X-rays. The detected backscattered electrons create a backscattered electron image (BEI) which shows a view depending on the composition of the sample. By detecting the X-ray emission the chemical composition of the sample can be obtained. [16] The SEM used in this study is a JSM-5800 produced by Jeol. The microscope is equipped with an electron dispersive spectroscopy (EDS) system, Link ISIS and a Link EDX (electron dispersive X-ray) detector, which are used for the chemical analyses. The entire equipment is shown in Figure 9. The SEM can create images with a resolution of 3.5 nm and high quality images can be obtained at a magnification of up to 10 000x. The EDS system can detect elements with atomic number 5 (boron) and higher.

Figure 8: Schematic image of SEM [15].

18

Figure 9: Jeol JSM-5800 Scanning Electron Microscope.

1.6.4.5 Light optical microscope

The light optical microscope (LOM) Olympus BX60M was used in the microstructural examination. The microscope is equipped with a video camera of the type Olympus U-TV1X which is connected to a computer. With this setup, shown in Figure 10, it’s possible to take pictures with a magnitude of up to 500x. The microscope has been used for taking overview photos and for examination of the carbides in Haynes 230 by means of the image analysis software Image-Pro Plus.

Figure 10: Olympus BX60M light optical microscope.

19

1.6.4.6 Stereo microscope

In order to examine the fracture surfaces of the tensile test specimens and to measure the angle of the slip planes in CMSX-4, the stereo microscope Leica Wild M10 was employed. The microscope is connected to a digital camera of the type Leica DC180 and the taken photos are transferred to a computer.

1.6.4.7 Hardness test equipment

Two different types of hardness testers have been employed in this study. The first one, a Diatestor 2N from Wolpert (see Figure 11), was used to determine the macrohardness of the materials. It has a loading capacity of 1-250 kg. The other one was a Leco M-400DT microhardness tester (see Figure 12) connected to a Panasonic CCTV camera. The photos were transferred to a computer where the measuring of the imprints was made. The microhardness tester has a loading capacity of 10-1000 g.

Figure 11: Macrohardness tester.

Figure 12: Microhardness tester.

20

2 SUPERALLOYS

The development of gas turbines has been the main driving force for the existence of the superalloys. The metallic alloys are today used in high-temperature and corrosion-resistant applications in a variety of industries. Not infrequently the service temperature goes beyond 0.7 of the superalloy’s melting temperature. Superalloys are divided into three classes: nickel-based, cobalt-based and iron-nickel-based alloys. The nickel-based superalloys are the superalloys most frequently used in gas turbine components. [17]-[20]

2.1 NICKEL-BASED SUPERALLOYS

Nickel-based superalloys can be used to a higher fraction of melting temperature and are therefore more favourable than cobalt-based and iron-nickel-based superalloys at service temperatures close to the melting temperature of the materials [18]. The application areas of nickel-based superalloys in gas turbines are for instance the turbine blades and vanes because of the alloys’ good corrosion and creep resistance and retained strength at high temperatures. Furthermore, the alloys are possible to strengthen in different ways and they are in possession of high phase stability. More about this is found in chapter 2.1.3 and 2.1.2, respectively.

2.1.1 Composition

The composition of nickel-based superalloys is altered depending on the desired properties. Besides nickel, the alloys contain in general 10-20 % chromium, up to 8 % aluminium together with titanium, and 5-10 % cobalt. Small amounts of boron, zirconium and carbon are included as well. Common addition in some alloys is for example molybdenum, tungsten, niobium, tantalum and hafnium. There are also some tramp elements, i.e. elements which unintentionally were included in the alloy, and these elements have to be carefully controlled. Examples of elements belonging to this group are silicon, phosphorus, sulphur, oxygen and nitrogen. [17] Chromium and aluminium are desired since they improve the oxidation resistance of the alloy. A small amount of yttrium binds the oxide layer to the substrate. Boron and zirconium are added to the polycrystalline superalloys where they segregate to the grain boundaries. This results in a better creep strength and ductility. The carbides tend to precipitate at the grain boundaries and prevent the sliding phenomenon of the boundaries. A few examples of carbide formers are carbon, chromium, molybdenum and tungsten. Some elements function as solid-solution strengtheners, e.g. cobalt, iron, niobium, rhenium and molybdenum. The addition of titanium will increase the hot corrosion resistance and the role of nickel is to give phase stability. [18], [20]

2.1.2 Phases and microstructure

Different phases are formed during the fabrication of nickel-based superalloys. The dominating phases are the gamma matrix and the gamma prime.

21

2.1.2.1 Gamma matrix (γ) phase

The foundation of all nickel-based superalloys is the gamma matrix, γ. The continuous phase is nonmagnetic and with a face centred cubic, FCC, structure [17]-[19]. Illustration is found in Figure 13. Nickel by itself has neither exceptional high elastic modulus nor low diffusivity. On the other hand nickel, has a nearly filled third electronic shell which allows alloying with solid solution strengthening elements without losses in phase stability. The alloy elements composing the gamma matrix mainly belong to Group V, VI and VII and are cobalt, iron, chromium, molybdenum and tungsten. [17]-[19]

2.1.2.2 Gamma prime (γ´) phase

The precipitated phase gamma prime, γ´, wasn’t identified until the 1940’s [12]. It is formed from elements from Group III, IV and V. The addition of for example aluminium and titanium, which are the essential solutes [20], results in a reaction with nickel, precipitating the γ´ phase of the form Ni3X, where X is an alloy element. The structure of the gamma prime is FCC. The atom ordering is illustrated in Figure 14. Other elements included in the γ´ phase can be chromium, hafnium, niobium and tantalum. [12], [17]-[20] The γ´ lattice parameter differs slightly from the one of gamma matrix. The mismatch is small, <0.2%, for spherical γ´. The close match makes it possible for the γ´ to precipitate evenly in the matrix. [12], [19] There are several benefits of the presence of γ´ in the matrix. The coherence between γ´ and γ results in a low surface energy and in an exceptional long-time stability. The phase is also the reason for the high-temperature strength and creep resistance in most superalloys. [17], [19] The strength of the alloy is strongly dependent on the volume fraction of γ´. Wrought alloys contain 20-45 % of γ´. Higher fractions will make the deformation too difficult. Cast superalloys can have a volume fraction of up to 60 %, which will increase the alloy strength compared with the wrought alloys. [19]

Figure 13: FCC gamma structure [20].

Figure 14: FCC gamma prime structure [20].

22

2.1.2.3 Carbides

The addition of carbon, generally in amounts of about 0.02 to 0.2 %, will result in a union with refractory elements such as hafnium, niobium and titanium. Primary carbides of the type MC, where M can be one of the just previously mentioned metals, are formed during freezing of the alloy. The structure of these carbides is FCC. MC carbides are an essential source of carbon during heat treatment and service. On those occasions, the carbides tend to decompose into other, secondary, carbide variants, e.g. M23C6 and M6C. The dominating formulas for formation of the two carbides are, according to Sims et al. [17], believed to be

'623 γγ +→+ CMMC (1)

and

'6 γγ +→+ CMMC . (2) In some alloys the formation of M6C from M23C6, according to the formula

MCMMCM ′′+→′+ 6236 , (3)

has been observed, and in yet other alloys the reverse reaction occurs. M´ and M´´ can be replaced with chromium and cobalt, nickel or molybdenum, respectively. [17] The most common M element in M23C6 is chromium but also iron, tungsten and molybdenum can be found in that position. The M elements generally found in M6C are molybdenum and tungsten, but the carbide can contain chromium, cobalt and tantalum as well. [19] The existence of carbides plays an important role in polycrystalline superalloys. Both M6C and M23C6 carbides are most likely to precipitate at grain boundaries. When properly formed, they strengthen the boundary and restrain grain boundary sliding [18]-[19]. When M23C6 is formed in the grain boundaries, the chromium content in the matrix is reduced and the solubility for γ´ is augmented in these zones [17]. The shape of the carbides is crucial for the properties. Cellular shape of the M23C6 carbides can cause premature failure while irregular, blocky particles will strengthen the alloy [19]. Fine, intergranular carbides will have a strengthening influence on the material. Deleterious elements in the superalloy can be tied up by the carbides; hence avoiding phase instability during service. [19]

2.1.2.4 Grain boundary (γ´) phase

It has been reported [17]-[18] that, in some alloys, γ´ may segregate to the grain boundaries when exposed to heat, forming a film around the M23C6 carbides or

23

the decomposing MC carbides. The existence of these films is believed to improve the rupture properties.

2.1.2.5 Borides

Small additions of boron have proved to have a positive influence on the creep-rupture resistance. The boron precipitates into the form M3B2 which has a tetragonal unit cell. Two different types of borides has been observed, where the M represented a mixture of cobalt, chromium, molybdenum, nickel and titanium. The borides are merely observed in grain boundaries. [17]-[19]

2.1.2.6 Topologically close-packed (TCP)-type phases

Not all phases found in the nickel-based superalloys are improving the properties. The composition has to be carefully controlled in order to avoid topologically close-packed (TCP) phases, for example σ phase, μ phase or Laves. These phases can be formed under certain conditions, usually during service. They are distinguished by their platelike or needlelike shapes. Alloys containing transition metals, such as tantalum, niobium, chromium, tungsten or molybdenum, are the alloys the most vulnerable to the formation of TCP phases. [17], [19] σ phase is a physically hard, platelike phase of the form (Fe,Mo)x(Ni,Co)y which has a seriously harmful influence on the properties of the alloy. Crack initiation caused by the shape and the hardness of the phase leads to brittle failure. Even more crucial is the effect on the strength of the alloy. The γ phase is depleted of refractory metals, resulting in a considerable loss of solution strengthening. The σ phase is also responsible for shortened rupture life, since high temperature rupture can occur along the plates. [17], [19] Platelike μ has been observed [19] but the effects of the phase on superalloys remain to be investigated. High content of Laves in the superalloy will decrease the tensile ductility and creep properties at room temperature. In the new generation superalloys the content of chromium is considerably reduced, since too much chromium tends to form TCP and hence deteriorate the properties of the alloy [20]. The subsequent reduction in corrosion protection is compensated by coating the component in question.

2.1.3 Strengthening of superalloys

Strengthening of superalloys is required for the purpose of obtaining the desirable high-temperature properties. It can be realised through either solid-solution strengthening or precipitation hardening. Creep resistance is an example of interaction between different hardening mechanisms. In early stages of creep the largest contributor to the creep resistance is the effects from solid-solution strengthening. The effect diminishes with time whereas the contribution from the precipitation hardening increases. [19]

2.1.3.1 Solid-solution strengthening

Solid solution is best described as a homogeneous crystalline structure in which one or more types of atoms or molecules may be partly substituted for the

24

original atoms or molecules without changing the structure. This substitution has a strengthening effect on the material. Common strengthening elements are chromium, cobalt, iron, molybdenum, rhenium, tantalum and tungsten. [19] As already mentioned in chapter 2.1.2.1, solid-solution strengthening takes place in γ phase. The addition of for example molybdenum expands the lattice and cobalt reduces the lattice when replacing iron in the superalloy matrix. An expansion of the lattice creates an internal strain. The expansion affects the mismatch with the strengthening precipitate phase. It has been shown that the stacking fault energy is reduced in the presence of solid-solution strengtheners. This will make it more difficult for dislocations, thus cross-slip at high temperatures is restrained. [17]-[18] Solid solution strengtheners can also have a beneficial influence on the corrosion and oxidation resistance. Superalloys with sufficiently large chromium content will form a protective oxide film that covers the surface [17].

2.1.3.2 Precipitation hardening

As described in chapter 2.1.2.2, gamma prime is formed during ageing by the precipitation of aluminium and titanium. The slightly different lattice parameter of γ´ creates a small misfit important for two reasons. First of all it guarantees a low γ/γ´ surface energy which is essential for a stable microstructure and improves the properties at elevated temperatures. Secondly, a negative misfit, i.e. γ´ has a smaller lattice parameter than γ, will facilitate the formation of rafts and by those means possibly reduce the creep rate. This will be discussed further in chapter 2.1.4. The misfit is controlled by the composition of the superalloy, particularly by altering the aluminium-titanium ratio, but also by the ageing temperature. [11], [17], [20]

2.1.4 Rafting

Appropriate heat treatment or service exposure can stimulate the formation of rafts. A negative misfit between γ and γ´ will result in internal stresses. Under this stress, the γ´ particles coalesce, forming layers in a direction perpendicular to the applied stress. These layers are also called rafts. [20]-[21] In low-stress applications, the rafts improve the creep resistance by preventing the dislocations to move over the layers. On the other hand, if the stress is large enough the dislocations will be able to cut through the rafts which have coarsened during formation. In this case, the presence of rafts is deleterious.

Figure 15: Rafts in single crystal superalloy [21].

25

2.2 GAS TURBINE BLADE CASTING

Gas turbine blades are exposed to severe conditions during service, such as elevated temperatures, high centrifugal forces and oxidation. Exceptionally good high-temperature properties are indispensable for the blade and a complex design is required to match these demands. Conventional casting will result in a superalloy with insufficient high-temperature properties caused by the structure containing equiaxed grains. Previously, gas turbine components made of conventionally cast superalloys failed at the grain boundaries from a combination of thermal fatigue, oxidation and creep. [19] The two methods used today, directionally solidified (DS) and single crystal (SX) cast superalloys, are superior to conventionally cast superalloys for two reasons. Grain boundaries are failure initiation sites and the alignment or elimination of these allows the γ´ to, through solution heat treatment, increase the creep strength and ductility at elevated temperatures of the alloy. Secondly, a low-modulus <001> orientation is created parallel to the solidification direction during directionally solidification, which improves the thermal fatigue resistance. [17], [19] Casting turbine blades is a complicated process because of the complex shape. Simplified, a mould is created by pouring a ceramic around a wax model of the component. The wax is then removed and the mould is filled with molten metal from the top of the blade. [20]

2.2.1 Directional solidification (DS)

The directional solidification of superalloys was not introduced until the early sixties. In the directional solidification process, the grain boundaries are aligned parallel to the solidification direction, which coincide with the principal stress axis of the component. The final structure consist of columnar grains with their <001> orientation in the direction of the later on applied load. [17] As mentioned before in chapter 2.2, stresses at elevated temperatures have a crucial effect on the grain boundary perpendicular to the stress direction, which is the weakest link of the chain. By aligning the grain boundaries the site of failure initiation is removed and the influence from the stress on the superalloy is very much reduced. [17] DS turbine blades with the low modulus <001> orientation parallel to the direction of the applied stress have proved [19] to increase the thermal fatigue resistance fivefold compared to conventionally cast turbine blades. A picture of a directionally solidified turbine blade is found in Figure 16. Directionally solidified components are not employed in any of the turbines produced by SIT AB.

Figure 16: Directionally

solidified columnar grains

turbine blade [20].

26

2.2.2 Single crystal (SX) casting

A small modification of the DS process, originally made in the mid-1960s, enables the casting of single crystal components. The principle is illustrated in Figure 17. The process is performed in vacuum in a preheated, ceramic mould.

Figure 17: Single crystal processing [19].

By controlling the solidification in the helical mould, the growth of all grains can be prevented; hence the presence of grain boundaries is eliminated. As shown in Figure 17b, the mould is chilled from the bottom. At the beginning the solidification results in columnar grains perpendicular to the thermal gradient. When reaching the helical channel, all columnar grains except one are prevented from growing. At the end of the helix, the structure of the solidifying superalloy is single crystal. [17], [19] A single crystal turbine blade with the helical channel is shown in Figure 18. Just as in the case of DS superalloys, the preferential orientation parallel to the stress axis is <001>. One reason to why single crystal superalloys exhibit better properties at high temperatures than DS superalloys is the lack of grain boundary strengthening solutes, i.e. boron and zirconium. The act of removing these elements will increase the incipient melting temperature of the superalloy thus improving the elevated-temperature properties. [20]

The use of single crystal material in the gas turbines produced by SIT AB is restricted to one single component: the first stage turbine blade in SGT-800. The remaining components are conventionally cast.

2.3 SINGLE CRYSTAL SUPERALLOY CMSX-4

CMSX-4 is a second generation single crystal superalloy from Cannon-Muskegon Corporation. The superalloy is characterised by a high γ´ solvus temperature and

Figure 18: Single crystal blade [20].

27

a high concentration of refractory elements [12]. Due to its high-temperature strength, oxidation resistance and creep resistance, it’s applied to the hottest parts of the gas turbine, for example the first stage turbine blade in SGT-800. CMSX-4 is developed from the first generation superalloys CMSX-2 and CMSX-3. The tantalum/tungsten ratio is balanced to improve the castability. The addition of 3 % rhenium improves both creep strength and thermal fatigue resistance of CMSX-4 compared to CMSX-2 and CMSX-3 at elevated temperatures. The added rhenium is mainly found in γ where it retards coarsening of γ´, thus increases the γ/γ´ misfit. Another change is the higher volume content of γ´ compared to the content in CMSX-2. The cause is to further improve the creep resistance. The solid solubility is assisted by the increased cobalt content and the reduced chromium content. Besides, this variation in composition prevents the formation of TCP at elevated temperatures. [19] CMSX-4 melts in the temperature range 1320-1380°C [12]. The material used in this study comes from two different batches, V 8508 and V8677, provided by Cannon Muskegon Corporation. The composition of CMSX-4 is shown in Table 1, values expressed in wt%. The mixture is balanced (Bal) with nickel.

Al Co Cr Hf Mo V 8508 5.66 9.7 6.4 0.10 0.61 V 8677 5.66 9.6 6.4 0.10 0.60

Ni Re Ta Ti W V 8508 Bal 2.9 6.5 1.04 6.4 V 8677 Bal 2.9 6.5 1.04 6.4

Table 1: Composition of CMSX-4 used in this study, expressed in wt% [22].

The superalloy is heat treated in three different stages. Firstly the material is heat treated under vacuum. The heat treatment commences at 1277°C ± 5°C and is in 7 steps increased to a temperature of 1321°C ± 3°C during 18 h in total followed by a rapid gas fan quench in argon. Secondly the material is held under vacuum at a temperature of 1140°C ± 10°C during 2 h and is then air cooled. Finally the material is held under vacuum at a temperature of 870°C ± 5°C for 20 h and is then air cooled. [22] The known, principal phases in CMSX-4 are gamma matrix and gamma prime. The fraction of gamma prime amounts to approximately 60 vol%. The possible formation of TCP at elevated temperatures is an unwanted occurrence to a large extent impeded by the addition of refractory elements. [11]

2.4 SUPERALLOY HAYNES 230

Haynes 230 is a nickel-based solid-solution strengthened superalloy with excellent creep-rupture strength at high temperatures and oxidation resistance. The superalloy is designed to sustain long-term thermal stresses without suffering from property degradation or grain coarsening. The melting

28

temperature range is 1301-1371°C and the maximum service temperature is 1149°C, i.e. >0.8·Tm. [13] Nickel-based superalloys that are solid-solution strengthened through a significant addition of molybdenum or tungsten, such as Haynes 230, are harder to form than alloys lacking these elements [18]. The tungsten was chosen as the primary solid-solution strengthener because of its beneficial effect on the creep strength and its ability of decreasing the stacking fault energy [23]. The carbon quantity is kept low in the material since a content exceeding 0.15 % can result in a significantly diminished ductility due to carbide precipitation [18]. The resistance to oxidation is improved due to the addition of chromium to the alloy’s composition. [23]. The exact composition of Haynes 230 used in this study is described in Table 2. The material is provided and analysed by Haynes International. The whole of the test certificate is shown in Annex 8.2.

Al B C Co Cr Cu Fe La 0.37 0.003 0.10 0.25 22.00 0.02 1.45 0.016

Mn Mo Ni P S Si Ti W 0.54 1.29 Bal <0.005 <0.004 0.44 <0.01 14.32

Table 2: Composition of Haynes 230 used in this study, expressed in wt% [Annex 8.2].

The superalloy can be either cast or wrought, depending on the application. Haynes 230 is produced in a variety of forms, e.g. plates, sheets and bars [13]. In order to form the sheets used in the combustion chamber, the superalloy is first cast and then rolled. However, the test specimens in this study was turned from a wrought bar. The heat treatment is carried out at a temperature between 1177°C and 1246°C and is then followed by rapid cooling [13]. The wrought superalloy consists mainly of the matrix and primary MC carbides [17]. During long-term exposure to heat, secondary carbides and intermetallic compounds can precipitate, resulting in for example beneficial creep strength improvement or reduced fracture resistance, depending on the precipitated phase [23].

29

3 AGEING EFFECTS IN SUPERALLOYS

3.1 CONTRACTION IN SUPERALLOYS

Shrinkage during service of superalloys and the underlying mechanisms have been discussed in several studies [1]-[2], [24]-[27]. Marucco [24] measured lattice contraction after ageing at intermediate temperatures (450-600°C) in 20Cr-25Ni steel, Sanicro 71 and Inconel 690 which amounted to 0.024, 0.040 and 0.038 % respectively. The study did also show that small ordered areas arose in the alloys due to the ageing. Nath et al. have been involved in several studies [2], [25]-[26] investigating the ageing effects on different nickel-based alloys, e.g. Nimonic 80A [26], a wrought nickel-based precipitation-hardened superalloy used as a bolting material in steam turbines. By use of X-ray diffractometry they proved that the lattice parameter was reduced due to ageing at 450-565°C and the kinetics of the contraction was mapped. The maximum lattice parameter contraction was obtained during ageing at 450°C for 30 000 h and measured 0.115 %. According to Nath et al. the lattice contraction was due to short range ordering (SRO) and long range ordering (LRO) arisen during ageing. Electron diffraction studies performed on Nimonic 80A corroborated the existence of SRO based on Ni2Cr and the transformation into LRO owing to long-term ageing. The lattice parameter contraction depends on the composition of the ordered phase. Marucco and Nath [2] have measured the lattice parameter contraction in Ni2Cr and Ni3Cr after ageing at 475°C (see Figure 19). Ni3Cr exhibited a small contraction in the neighbourhood of 0.05 % after 10 000 h. The lattice parameter contraction in Ni2Cr under the same conditions is almost 5 times larger. In previous work carried out at SIT AB [1] the lattice contraction in different nickel-based superalloys and the austenitic stainless steel X6CrNiTi18-10 was measured. Cylindrical specimens of the length 100 mm were aged at 450, 500 and 550°C and the shrinkage strain was measured after 300, 1000 and 3000 hours in a coordinate measurement machine. The result is shown in Figure 20-Figure 22.

Figure 19: Lattice contraction in (○) Ni2Cr and (□) Ni3Cr as a function of the ageing time at 475°C [2].

30

Shrinkage of Austenitic Steels and Ni-base alloys Temperature: 450degC

-0,06-0,054-0,048-0,042-0,036-0,03

-0,024-0,018-0,012-0,006

0

0 500 1000 1500 2000 2500 3000 3500Time, h

Stra

in, %

Hastelloy XHaynes 230Inconel 718Haynes HR-120233702

Figure 20: Shrinkage strain at 450°C [1].


-0,06

-0,05

-0,04

-0,03

-0,02

-0,01

0

0 500 1000 1500 2000 2500 3000 3500Time, h

Stra

in, %




-0,06

-0,05

-0,04

-0,03

-0,02

-0,01

0

0 500 1000 1500 2000 2500 3000 3500Time, h

Stra

in, %



31

It is obvious that the contraction of the alloys is significant already after 300 h of ageing. Between 300 h and 1000 h the shrinkage rate is slowing down and is almost constant between 1000 h and 3000 h. This behaviour is in accordance with the conclusion Nath et al. made in 1984 [26] which is further discussed above Figure 25 in chapter 3.2. The shrinkage strain after 3000 h of ageing is plotted against the ageing temperature in Figure 23 and the value of the maximum shrinkage strain is presented in Table 3.

Shrinkage strain after 3000h ageing

0

100

200

300

400

500

600

-0,06 -0,04 -0,02 0Strain, %

Tem

pera

ture

, C


Figure 23: C-curves of shrinkage strain at ageing at 3000 h [1].

Material Maximum shrinkage strain [%]

Ageing temperature

[°C]

Hastelloy X -0.036 450 Haynes 230 -0.052 500 Inconel 718 -0.02 450

Haynes HR-120 -0.016 450 X6CrNiTi18-10 -0.024 550

Table 3: Maximum shrinkage strain for different austenitic alloys after ageing for 3000 h [1].

The largest shrinkage is attributed to Haynes 230 at ageing at 500°C and then Hastelloy X at ageing at 450°C. The lattice parameter of the alloys was measured with X-ray diffraction before and after ageing. A good correlation was found between macroscopic length shrinkage and lattice parameter contraction.

32

3.2 SHORT RANGE ORDERING (SRO) AND LONG RANGE ORDERING (LRO)

Short range ordering exists in all Ni-Cr alloys, irrespective of the presence of other alloy elements. Small, ordered areas in the size of nanometres are formed in an otherwise disordered matrix. The ordering takes place if atoms of different types are more attracted to each other than atoms of the same type. The short range ordered phase forms at stoichiometric compositions, e.g. AB, A2B and A3B or at off-stoichiometric compositions but at lower kinetics. A commonly formed ordering phase in superalloys is the orthorhombic Ni2Cr. The formation takes place during either cooling from solution temperature or at early stages of ageing. The degree of SRO decreases with temperature, but it is also influenced by the composition. Studies on ternary Ni-Cr-Fe alloys [28] have shown that the augmentation of Ni content increases the degree of SRO and the diminution of Fe content has the opposite effect. [2], [24], [28] Long-term ageing below a critical temperature TC will, for certain alloy compositions, result in growth of SRO nuclei, transforming the SRO phase into long range ordering phase [28]-[29]. TC depends on the composition of the alloy, but is normally located between 530°C and 580°C [28]. Even if LRO is non-existing above TC, SRO can exist at temperatures close above TC in stoichiometric alloys AxBy, alloys which would be long range ordered at lower temperatures [2]. See Figure 24 for illustration.

Figure 24: The dependence of the long (S) and short (σ) range

order parameters on temperature [2].

According to Rtishchev [29] SRO and LRO are formed in binary Ni-Cr alloys with a chromium content of 25-37 at%. In commercial superalloys, the chromium content in γ phase after the γ´ precipitation has to exceed 25 at% to make the formation of LRO possible. The tendency of Ni-Cr based superalloys to form LRO can be described by the Z-criterion [29]. Z is calculated on the atom content in γ, after precipitation of γ´ and minor phases, according to Equation 2:

33

MoWCrNiZ++

= . Equation 2: Z-criterion, in at%·at%-1 [29].

Rtishchev has proved experimentally that the formation of LRO occurs at estimated values of Z<3.0. The Ni2Cr superlattice stability is affected by the alloy elements. It has been shown that the presence of tungsten and molybdenum stabilizes the superlattice while the presence of cobalt has a negative effect on the stability of the structure [29]-[30]. The degree of LRO and the kinetics depend on the ageing temperature and the composition. If the ageing temperature is increased, the nucleation rate will decrease while the growth rate will increase. The maximum kinetics is obtained for the composition Ni2Cr. The larger deviation from the stoichiometry the slower kinetics of LRO is obtained [28]. The presence of tungsten and molybdenum will have an accelerating effect on the order kinetics. On the other hand, the kinetics is strongly subdued by the presence of cobalt and iron. [24], [28]-[29] Nath et al. [26] showed that the kinetics of lattice contraction depends on the ageing time and applied strain. The results from their lattice contraction measurements for Nimonic 80A are plotted in Figure 25. They concluded that the contraction is divided in three stages. During the first 500 h the contraction rate was high. During the next 15 000 h it became very slow and then it accelerated again at longer ageing time.

Figure 25: Lattice parameters of strained and unstrained Nimonic 80A during ageing: ○ – 500°C with no strain, ● – 500°C with applied strain, □ – 450°C with

no strain, ■ – 450°C with applied strain [26].

34

The lattice parameters at 450°C and 500°C were similar for the majority of the ageing interval. The exception is located in the early stages of ageing where the material with applied strain contracted faster at 450°C than at 500°C. Noteworthy is also that the unstressed alloys initially contracted faster than the stressed alloys, but after a specific ageing time the rate of lattice contraction of the unstressed alloys slowed down, resulting in a crossover of the curves. This crossover takes place after about 3000 h at 500°C and after 25 000 h at 450°C. The ordering phenomenon affects the properties of the superalloy, mainly through influence on the dislocation morphology. Ordering arise since the attraction is larger between atoms of different types than between atoms of the same type. Dislocation moving in a superalloy with such structure along one slip plane would result in a Ni atom taking the place of a Cr atom. In this case, a superdislocation is created. The superdislocation dissociates into two “normal” dislocations and one antiphase boundary. If the ordered domains are smaller than the width of the antiphase boundary, for example in the early stages of LRO, the dislocation moving along one particular slip plane preclude the nearest-neighbour bonding across the slip plane. To avoid increased disorder, successive dislocation moving must occur along the same slip plane. This results in an increased ease of dislocation to occur on a plane where slips already have taken place. The dislocation is restricted to a small number of slip planes where it takes place with coarse steps. This behaviour is also called heterogeneous deformation. [2], [31] The deformation is remained planar if the absorption of the dislocations, caused by high friction stresses [30], is as fast as the emergence at grain boundaries. Then there will be no work hardening of the alloy. If the annihilation process is too slow a pile-up will be created and further deformation on the plane is limited. This will result in an activation of other planes and the presence of dislocation tangles between these planes will cause a strain hardening in the alloy. [2] Rtishchev [29] found that the yield strength is significantly improved by the presence of LRO. Unfortunately this leads to sharp notch sensitivity at stress rupture and deteriorated ductility. The ordering is also responsible for an increase in elastic modulus as well as in hardness. Marucco and Nath [2] found that the tensile strength and ductility decreased with increasing test temperature and that the minimum work hardening coefficient was found between 450 and 600°C. They also found high rates of work hardening at room temperature, but at higher temperatures the dislocation-annihilation will be faster, thus encouraging planar slip. The ordering does also affect the electrical resistivity, ρ. In a study made of Marucco and Nath [2] the resistivity was measured during long-term ageing at 475°C (see Figure 26). A comparison between the ordered phases Ni2Cr and Ni3Cr shows a large difference in resistivity change. The resistivity in Ni3Cr alloys increased with approximately 0.4 % already after 24h. ρ was then more or less constant during ageing up to 25 000h. The structure of the alloys remained in the SRO state during the entire ageing time. The Ni2Cr alloys showed a similar microstructure and resistivity at early stages of ageing. However, after a few hundred hours the transformation of SRO into LRO begins, and the resistivity decreases. After long-term ageing decrease in resistivity of up to 60 % was measured.

35

Figure 26: Resistivity changes in (○) Ni2Cr and (□) Ni3Cr as a function of the ageing time at 475°C [2].

The order-disorder transformation takes place in simple commercial nickel-based alloys such as Nimonic 80A, but the LRO phase with Ni2Cr structure has been found in more complex superalloys as well, such as Hastelloy [2].

3.3 NEGATIVE CREEP

The term negative creep is used for the superposition of two opposite processes. Common plastic creep represents the positive component and lattice contraction due to ordering represents the negative component. Negative creep in superalloys can be observed as a contraction during creep test or as an increase in stress during relaxation test. [2], [27] It has been shown [2] that the ordering of Ni2Cr in the form SRO and LRO arisen during ageing at 550°C and below is responsible of lattice contraction, which leads to dimensional instability and negative creep in Nimonic 80A and the binary alloy Ni-20wt%Cr. It was also confirmed that the precipitation of γ´ had no influence on the process [2], [30].

3.4 CARBIDE PRECIPITATION

In an internal report from Lincoln [32] the ageing effect on macrohardness of Haynes 230 is investigated. The result is shown in Figure 27. The test load used in the study was 20 kg. The report is analysing the influence of ageing temperature (750-1050°C) and ageing time (1000, 3000 and 5000 h) on the hardness. As illustrated in the figure, ageing at lower temperatures resulted in an improved hardness, but the hardness decreased with increasing temperature and the higher ageing temperatures resulted in a deteriorated hardness compared to the unaged reference.

36

Figure 27: Hardness changes due to ageing. [32]

37

4 AGEING INFLUENCE ON MECHANICAL PROPERTIES AND MICROSTRUCTURE OF HAYNES 230 AT INTERMEDIATE TEMPERATURES (500-550°C)

4.1 SPECIMEN PREPARATION AND EXPERIMENTAL PROCEDURE

4.1.1 LCF test

Nine cylindrical LCF test specimens were turned from the same bar, all of them following the drawing 7000 3005-1 (see Annex 8.3.1). Six specimens were aged during 1000 h, three of them at 500°C and another three at 550°C. The remaining 3 specimens were kept untreated as references (ref). LCF tests were performed on the specimens in the universal testing machine Instron 1275 at room temperature in agreement with ASTM E 606-92. The nominal diameter of the specimens was 10 mm and the extensometer gauge length was 12.5 mm. The tests were carried out in total strain control mode. The strain rate was 6 %·min-1 and the nominal strain ranges applied in the tests were 1.2 %, 1.6 % and 2.0 %. At each strain range, one reference and specimens from 500°C and 550°C respectively were tested, i.e. in total three specimens at each strain range. The number of cycles in each LCF test, N, was counted during the test. The number of cycles to crack initiation, Ni, was defined as 5 % load drop from the stabilised loop and the number of cycles to failure, Nf, as 50 % load drop. In those tests where the crack initiation took place outside the extensometer range, Nf was defined as the cycle when complete failure occurred. The maximum stress was determined for the first cycle, σmax (N=1), and for midlife, σmax (Nf/2). The plastic strain range, Δεpl, was calculated for midlife. All details about the calculations are given in Annex 8.4.1.

4.1.2 Impact test

The impact tests were performed at room temperature in Charpy impact test equipment in agreement with the standard EN 10 045-1. The impact test specimens were milled from a round bar into the wanted shape according to the drawing 7000 2059-3 (see Annex 8.3.2). Ten of the fifteen V-notched specimens were aged for 1000 h, five of them at 500°C and the other five at 550°C. The last five which weren’t heat-treated represented the reference.

4.1.3 Tensile test

The specimens for the tensile test were turned from the same bar, following the drawing 7000 1782-3 (see Annex 8.3.3). Out of totally twenty specimens, four were used as references. Two were aged during 300 h, three during 1000 h and another three during 2500 h, all at 500°C. This procedure was repeated with another set of specimens at 550°C (see Table 4). To carry out the tensile test in room temperature, Schenck-Trebel RSA 100 test machine was used. All tests were performed within the standard EN 10 002-1. The strain was measured by an extensometer up to 5.0 %.

38

Specimen marking

Ageing temp [°C]

Ageingtime [h]

Specimen marking

Ageing temp [°C]

Ageing time [h]

HY21 500 1000 HY31 550 2500 HY22 500 1000 HY32 550 2500 HY23 500 1000 HY33 500 300 HY24 500 2500 HY34 500 300 HY25 500 2500 HY35 550 300 HY26 500 2500 HY36 550 300 HY27 550 1000 HY37 - - HY28 550 1000 HY38 - - HY29 550 1000 HY39 - - HY30 550 2500 HY40 - -

Table 4: Ageing temperature (temp) and time for Haynes 230 tensile test specimens.

The specimens’ dimensions were measured before and after the test and the elongation, A5, was calculated as well as the reduction of the cross-section area, Z and the yield strength at 0.1 % plastic strain deformation, Rp0.1. The tensile strength, Rm, and the yield strength at 0.2 % plastic strain deformation, Rp0.2, were determined by the employed software during the test. No conclusions were drawn from the elastic modulus, E, calculated by the software, since the determination of E from tensile test curves is difficult and the result can be unreliable. All details concerning the calculations are given in Annex 8.4.2. All specimens aged under the same conditions, time and temperature, form a group. Since each group contain several curves, a graph with all curves aged at for example 500°C would be undecipherable. In order to make the graphs clear, an attempt was made to use the Ramberg-Osgood equation

n

KE′

⎟⎠⎞

⎜⎝⎛

′+=

1

σσε Equation 3: Ramberg-Osgood equation.

to produce a fitted curve representing an average for each set of curves. In this expression K´ is the strength coefficient and n´ the strain-hardening exponent. However, the fitted curve matched up poorly to the original curves, so this adjustment was inapplicable. Instead, one curve from each group which best described the behaviour of the group was distinguished and plotted in Figure 34 and Figure 35, found in chapter 4.2.3.

4.1.4 Microscopic examination

The cut-ups prepared for microscopic examination contain a cross-section of the tensile test specimens or the 2.0 % strain range LCF test specimens along the test load axis. The chosen specimens were cut perpendicular to the test load axis using a slitting wheel. Approximately 3 cm of the specimens were cut off,

39

including the fracture surface, and the samples were then mounted. Cut-ups containing tensile specimens were grinded roughly on a stone down to the core of the specimens. LCF specimens were sectioned along the test load axis in the slitting machine before the mounting. All cut-ups were then ground, starting with 120 grade paper finish down to 1000 in five steps. The cut-ups were then polished with diamond suspension, first with the diamond size 6 μm and finally with the size 1 μm. The last act was electrolytic etching in 10 % oxalic acid at 6 V for 25-30 s. These etching conditions were excellent for the entire area of the specimens except for the edges. After microscopic examination of the cut-ups, the grinding and polishing procedure was repeated from 600 grade paper finish and completed with electrolytic etching for 5 s to enable examination of the edges. The fracture surfaces of the tensile specimens were photographed in the stereomicroscope before mounting. The cut-ups were examined in both optical microscope and SEM. The examination focused on any possible difference between new and aged material, such as carbide size and distribution, growth of secondary carbides, deformation, new phases and composition of both matrix and carbides. The carbide size was measured by means of the software Image-Pro Plus. The mean diameter, d1, and the aspect ratio, r1, were determined for primary carbides, and also the diameter and aspect ratio for secondary carbides, d2 and r2. Around 100 carbides were dimensioned in each sample. In the case where the carbides in the tensile test specimens were measured, the photos in question were taken approximately 3 cm from the fracture surface. When it comes to the untested specimens, the photos were taken in the middle of the sample.

4.1.5 Hardness test

Hardness test was performed on the cut-ups used in the microscopic examination before the re-etching. The macrohardness was measured in Vickers scale (HV10) in a macrohardness test machine from Wolpert Probat of the type Dia testor 2N. Four imprints were made on each specimen. The minimal distance between two imprints was at least three times the width of an imprint. The applied load was 10 kg. The test was carried out on the LCF test specimens as well as on a set of untested specimens. One of the untested specimens was a virgin bar and the other two were aged at 550 °C for 1000 h and 3000 h respectively. The size of each imprint was measured manually and the value of the Vickers hardness was then obtained from a table. Leco M-400 DT was employed to measure the microhardness of the untested specimens. The imprint had to be the largest possible to secure the reliability of the imprint measurement and yet small enough to enable the avoidance of the carbides. Experiments showed that 50 g was the optimal load. The distance between the imprints followed the same rules as for macrohardness testing. Imprints were made at two different areas: close to grain boundary and in the middle of the grain. Ten imprints were made in the matrix at each area and sample, avoiding the primary carbides. The size of the imprints was measured by means of the software Image Access Hardness and the value of Vickers hardness was calculated automatically.

40

4.2 TEST RESULTS

4.2.1 LCF test

Nine LCF tests were performed. The results from the LCF tests are summarised in Table 5.

Specimen marking

Ageing temp [°C]

Strain range

[%]

Ni Nf σmax (N=1) [MPa]

σmax(Nf/2) [MPa]

Δεpl(Nf/2)

L9 - 1.2 8093 8122 427 511 0.734 L5 500 1.2 7142 7193 430 538 0.718 L4 550 1.2 6748 6849 439 554 0.715

L13 - 1.6 2737 2771 435 563 1.058 L6 500 1.6 3035 3068 442 610 1.046 L8 550 1.6 3889 3889 428 585 1.042

L14 - 2.0 1783 1793 455 598 1.419 L7 500 2.0 1670 1684 446 654 1.387

L12 550 2.0 1982 1982 441 633 1.377

Table 5: Summary of result from LCF test.

Figure 28 and Figure 29 show the evolution of maximum stress and plastic strain range plotted versus the number of cycles.

4 0 0

4 5 0

5 0 0

5 5 0

6 0 0

6 5 0

7 0 0

7 5 0

1 1 0 1 0 0 1 0 0 0 1 0 0 0 0C y c l e

M a x s t r e s s[ M P a ]

n e w - 1 . 2 % 5 0 0 ° C - 1 . 2 % 5 5 0 ° C - 1 . 2 %n e w - 1 . 6 % 5 0 0 ° C - 1 . 6 % 5 5 0 ° C - 1 . 6 %n e w - 2 . 0 % 5 0 0 ° C - 2 . 0 % 5 5 0 ° C - 2 . 0 %

Figure 28: Max stress versus number of cycles.

41

0 , 6

0 , 8

1

1 , 2

1 , 4

1 , 6

1 1 0 1 0 0 1 0 0 0 1 0 0 0 0

C y c l e

P l a s t i cs t r a i n

r a n g e [ % ]

n e w - 1 . 2 % 5 0 0 ° C - 1 . 2 % 5 5 0 ° C - 1 . 2 %n e w - 1 . 6 % 5 0 0 ° C - 1 . 6 % 5 5 0 ° C - 1 . 6 %n e w - 2 . 0 % 5 0 0 ° C - 2 . 0 % 5 5 0 ° C - 2 . 0 %

Figure 29: Plastic strain range versus cycle.

The maximum stress, the inelastic strain and the total strain range at midlife is plotted against the number of cycles, one curve for each ageing condition. The curves are found in Figure 30-Figure 32.

M a x i m u m s t r e s s v e r s u s c y c l e a t m i d l i f e

4 5 0

5 0 0

5 5 0

6 0 0

6 5 0

7 0 0

1 0 0 1 0 0 0 1 0 0 0 0

C y c l e

σ m a x

[ M P a ]

r e f5 0 0 ° C5 5 0 ° C

Figure 30: Maximum stress versus cycle at midlife.

42

I n e l a s t i c s t r a i n v e r s u s c y c l e a t m i d l i f e

0 , 60 , 70 , 80 , 9

11 , 11 , 21 , 31 , 41 , 5

1 0 0 1 0 0 0 1 0 0 0 0

C y c l e

Δ ε p l

[ % ]

r e f5 0 0 ° C5 5 0 ° C

Figure 31: Inelastic strain versus cycle at midlife.

S t r a i n r a n g e v e r s u s c y c l e a t m i d l i f e

0 , 6

0 , 81

1 , 2

1 , 4

1 , 61 , 8

2

2 , 2

1 0 0 1 0 0 0 1 0 0 0 0

C y c l e

Δ ε t o t

[ % ]

r e f5 0 0 ° C5 5 0 ° C

Figure 32: Total strain range versus cycle at midlife.

4.2.2 Impact test

The impact test was carried out on 15 specimens, 5 from each group: new material, aged at 500°C and aged at 550°C, respectively. The result is shown in a chart in Figure 33 together with the values of the absorbed energy. Each column represents an average from the 5 specimens prepared under the same conditions.

43

1 3 9 , 2

1 2 1 , 01 0 6 , 3

0 , 0

4 0 , 0

8 0 , 0

1 2 0 , 0

1 6 0 , 0

2 0 0 , 0

n e w A g e d a t5 0 0 ° C

A g e d a t5 5 0 ° C

A b s o r b e d e n e r g y [ J ]

Figure 33: Result from impact test of Haynes 230.

The absorbed energy, KV, and the standard deviation are found in Table 6.

Ageing temp [°C]

KV [J]

Standard deviation

- 106.3 4.28 500 121.0 10.44 550 139.2 13.05

Table 6: Impact test results from specimens aged for 1000 h.

4.2.3 Tensile test

The total number of tensile tests performed was 20, of which one test was stopped prematurely by accident. Since the stop occurred close upon the expected failure, the test data was still of use, tensile strength excepted. The results from the tensile testing are shown in Table 7. All values are calculated averages from each group of specimens prepared under the same conditions, each group comprising 2-4 specimens. The results from the testing of the aged material are shown as the difference between new and aged material, written in percentages. The complete table with all original test results is found in Annex 8.5.1, Table 18.

44

Ref 300 h 1000 h 2500 h 500°C

[%] 550°C

[%] 500°C

[%] 550°C

[%] 500°C

[%] 550°C

[%]

Rm 862 MPa 1.0 0.9 1.0 1.3 1.7 1.2 Rp0.1 366 MPa 9.5 9.5 8.7 9.6 10.3 13.3 Rp0.2 379 MPa 4.0 3.2 3.9 4.7 5.5 5.6 A5 47.8 % 6.1 12.3 5.9 6.0 9.8 8.4 Z 48.0 % 16.2 20.0 13.2 4.6 16.4 16.3

Table 7: The percentage increase of tensile properties compared with the reference.

The stress-strain curves plotted up to 5 % strain are displayed in Figure 34 and Figure 35.

T e n s i l e t e s t , a g e d 5 0 0 ° C

0

1 0 0

2 0 0

3 0 0

4 0 0

5 0 0

6 0 0

0 1 2 3 4 5 6

S t r a i n [ % ]

S t r e s s[ M P a ]

H Y 4 0 - r e f H Y 3 3 - a g e d 3 0 0 hH Y 2 1 - a g e d 1 0 0 0 h H Y 2 6 - a g e d 2 5 0 0 h

Figure 34: Stress-strain curves from tensile test, ageing temperature 500°C.

45

T e n s i l e t e s t , a g e d 5 5 0 ° C

0

1 0 0

2 0 0

3 0 0

4 0 0

5 0 0

6 0 0

0 1 2 3 4 5 6

S t r a i n [ % ]

S t r e s s[ M P a ]

H Y 4 0 - r e f H Y 3 6 - a g e d 3 0 0 hH Y 2 8 - a g e d 1 0 0 0 h H Y 3 0 - a g e d 2 5 0 0 h

Figure 35: Stress-strain curves from tensile test, ageing temperature 550°C.

4.2.4 Hardness test

The results from the macrohardness testing are visualised in Table 8. The column of Vickers hardness contains the calculated averages from all imprints made on each sample. The last column shows the increase or decrease in Vickers hardness for the aged specimens compared with the unaged reference of the same type, i.e. untested or LCF tested, expressed in percentages.

Specimen type

Ageing temp [°C]

Ageingtime [h]

Vickers hardness[HV10]

Standarddeviation

Difference compared to ref [%]

Untested - - 202.8 4.03 Untested 550 1000 211.8 4.86 4.4 Untested 550 3000 219.3 4.65 8.1

LCF tested - - 280.0 5.23 LCF tested 500 1000 280.3 5.74 0.1 LCF tested 550 1000 275.0 7.35 -1.8

Table 8: Summary of results from macrohardness tests.

46

The results from the macrohardness testing for the, until then, untested specimens are shown in Figure 36 and the result from the previously LCF tested specimens are shown in Figure 37.

+ 4 . 4 %+ 8 . 1 %

1 8 5 ,0

1 9 0 ,0

1 9 5 ,0

2 0 0 ,0

2 0 5 ,0

2 1 0 ,0

2 1 5 ,0

2 2 0 ,0

2 2 5 ,0

2 3 0 ,0

0 1 0 0 0 3 0 0 0

a g e i n g t i m e [ h ]

V i c k e r sh a r d n e s s

[ H V 1 0 ]

Figure 36: Results from macrohardness test, untested

specimens aged at 550°C.

- 1 .8 %+ 0 .1 %

2 5 0 , 0

2 5 5 , 0

2 6 0 , 0

2 6 5 , 0

2 7 0 , 0

2 7 5 , 0

2 8 0 , 0

2 8 5 , 0

2 9 0 , 0

0 5 0 0 5 5 0

a g e i n g t e m p [ ° C ]

V i c k e r sh a r d n e s s

[ H V 1 0 ]

Figure 37: Results from macrohardness test, LCF specimens

aged for 1000 h.

The microhardness testing was only performed on the untested samples. The results are shown in Table 9 and in a chart in Figure 38. The given values of Vickers hardness is the calculated average from the ten imprints on each sample.

47

Specimen type

Ageing time [h]

Tested area

Vickers hardness[HV0.05]

Standard deviation


Untested - Close to grain boundary

222.3 12.07

Untested 1000 Close to grain boundary

264.1 9.49 18.8

Untested 3000 Close to grain boundary

244.6 12.63 10.0

Untested - Middle of the grain

228.9 13.04

Untested 1000 Middle of the grain

252.4 12.14 10.3

Untested 3000 Middle of the grain

236.7 10.23 3.4

Table 9: Summary of results from microhardness test, specimens aged at 550°C.

+ 1 8 . 8 %

+ 1 0 . 0 %+ 1 0 . 3 %

+ 3 . 4 %

2 0 0 , 0

2 1 5 , 0

2 3 0 , 0

2 4 5 , 0

2 6 0 , 0

2 7 5 , 0

2 9 0 , 0

3 0 5 , 0

0 1 0 0 0 3 0 0 0

A g e i n g t i m e [ h ]

V i c k e r sh a r d n e s s [ H V 0 . 0 5 ]

n e a r g r a i n b o u n d a r y m i d d l e o f t h e g r a i n

Figure 38: Results from microhardness test, untested specimens aged at 550°C.

4.2.5 Carbide size measurements

The estimated aspect ratio and mean diameter of the primary carbides are listed in Table 10 together with their min and max values and also the standard deviation, std. dev. The determination of the secondary carbides’ dimensions turned out to be very hazardous because of the inadequate photo resolution and limitations in the software. These results are therefore not included in the report.

48

Test type

Ageing time [h]

Ageing temp [°C]

r1 Std. dev.

Min Max d1 [μm]

Std. dev.

Min [μm]

Max [μm]

- - 1.71 0.64 1.02 4.51 1.84 1.86 0.20 9.05 300 500 1.50 0.40 1.00 3.20 2.20 1.69 0.28 8.60 300 550 2.55 1.97 1.00 10.29 2.06 2.04 0.28 8.90

1000 500 1.39 0.36 1.00 2.96 2.51 1.93 0.28 9.20

Tensile test

1000 550 1.89 1.20 1.00 7.57 1.62 1.55 0.28 7.39

- - 1.49 0.45 1.00 3.57 2.72 1.94 0.28 10.241000 550 1.58 0.73 1.01 4.76 2.55 1.45 0.34 8.38 Untested

3000 550 1.46 0.61 1.04 5.51 2.62 1.75 0.34 11.38

Table 10: Dimensions of primary carbides in Haynes 230.

An example of a LOM image used to measure the carbide size is found in Figure 39.

Figure 39: LOM image of carbides in material aged

300 h at 550°C.

4.2.6 Chemical analysis

The results from the chemical analysis are shown below. Several analyses were carried out for each specimen and area. The results presented in the following tables are representing the average result from each specimen and specific testing area. The analysed areas are the matrix (Table 11), primary carbides (Table 12) and secondary carbides (Table 13).

49

Ageing time [h]

Ageing temp. [°C]

C Al Ti Cr Mn Fe Co Ni Mo W

- - - 0.4 - 22.0 0.5 1.3 0.3 60.1 1.2 14.2 300 500 - 0.4 - 22.0 0.5 1.3 0.3 60.6 1.4 13.7 300 550 2.2 0.5 0.1 21.5 0.4 1.3 0.3 58.9 1.2 13.6

1000 500 - 0.3 - 22.0 0.6 1.3 0.3 60.2 1.3 13.9 1000 550 - 0.3 - 22.0 0.5 1.4 0.4 60.6 1.2 13.7

- - - 0.4 - 22.1 0.5 1.4 0.4 60.1 1.2 13.9 1000 550 - 0.4 - 22.5 0.6 1.4 - 61.0 1.0 13.0 3000 550 - 0.4 - 22.0 0.5 1.3 0.3 60.5 1.3 13.7

Table 11: Chemical analysis of the matrix in different specimens. The first five specimens are from the tensile test and the last three specimens are untested.

Ageing time [h]

Ageing temp. [°C]

Al Cr Mn Fe Co Ni Mo W

- - 0.2 13.5 - 0.3 0.2 20.5 4.3 60.9 300 500 0.2 13.5 - 0.5 - 20.3 4.0 61.6 300 550 0.3 13.5 - 0.4 - 20.0 4.3 61.5

1000 500 0.1 13.4 0.2 0.4 0.4 20.0 4.3 61.1 1000 550 0.2 13.5 - 0.4 0.2 20.8 4.1 60.8

- - 0.4 13.7 - 0.4 0.2 21.0 3.9 60.4 1000 550 0.4 13.5 - 0.5 0.2 20.2 4.1 61.1 3000 550 0.3 13.5 - 0.5 - 20.3 4.4 61.0

Table 12: Chemical analysis of primary carbides in different specimens. The first five specimens are from the tensile test and the last three specimens are untested.

Ageing time [h]

Ageing temp. [°C]

C Al Ti Cr Mn Fe Co Ni Mo W

- - - 0.2 0.1 24.5 0.7 1.4 0.4 66.0 0.6 6.2 300 500 - 0.2 - 23.1 0.5 1.4 0.3 65.0 0.7 8.7 300 550 - 0.5 - 21.7 0.5 1.4 0.3 59.3 1.3 15.1

1000 500 0.6 0.2 0.1 23.0 0.5 1.3 0.3 65.7 0.7 7.5 1000 550 7.0 0.5 0.2 20.8 0.4 1.2 0.4 52.8 1.5 15.3

- - - 0.4 - 21.9 0.6 1.3 0.4 59.4 1.3 14.7 1000 550 0.4 0.5 - 22.8 0.5 1.3 0.3 57.0 1.5 15.6 3000 550 2.4 0.5 - 22.5 0.4 1.2 0.3 53.7 1.7 17.6

Table 13: Chemical analysis of secondary carbides from grain boundaries in different specimens. The first five specimens are from the tensile test and the last three specimens

are untested.

50

4.3 DISCUSSION

4.3.1 Mechanical tests

The ageing has an interesting effect on the LCF properties. As shown in Table 5 and Figure 28-Figure 29, the maximum stress at midlife is visually affected by the ageing. At the lowest strain range, 1.2 %, the maximum stress increases with the ageing temperature, but at 1.6 and 2.0 % strain range the maximum stress is larger after ageing at 500°C than after ageing at 550°C. However the stress is still lower in the unaged material. The better illustration of the ageing influence is found in Figure 30. The graph clearly shows that the maximum stress is increased due to ageing. On the other hand, the total strain and inelastic strain is unaffected by the ageing. Figure 31 and Figure 32 show that the LCF properties for the aged specimens are equal to the behaviour of the unaged reference. The LCF properties are favoured by a fine grain size [23]. The grain size in Haynes 230 is controlled by the existence of primary carbides since the carbides resist solution during heat-treatment. By controlling the grain size the degradation of LCF properties due to ageing is prevented. As expected, the lifetime of the specimens, i.e. the number of cycles until failure, will shorten when the strain range is increased. However the lifetime is not affected appreciably by the ageing. The lifetime of the specimens exposed to the same nominal strain range is approximately the same, whether the material has been aged or not. Ageing of Haynes 230 does not only increase the resistance to plastic deformation, it increases the ductility as well. The experiments show that the impact resistance increases with increasing ageing temperature (see chapter 4.2.2). The absorbed energy increased with 14 % after ageing at 500°C for 1000 h and with 31 % after ageing at 550°C for 1000 h. The results from the tensile test (see Table 7) show that the ageing has a significant influence on the tensile properties. The yield strength is improved for specimens aged at both 500 and 550°C. The yield strength at 0.1 % plastic strain is around 10 % larger than the reference for all ageing times and temperatures. At 0.2 % plastic strain the yield strength has increased about 5 % compared to the unaged material. The tensile strength increases after ageing as well, by 1-2 %. The increase in tensile strength and yield strength is in good correlation with the results from the LCF test, where it was shown that the ageing results in an increased maximum stress in the material. The ageing time doesn’t seem to have any affect on the tensile properties; the results are more or less the same for each group of specimens aged at 300, 1000 and 2500 h. However, an interesting remark is the knee formation on the stress-strain curves representing the aged materials (see Figure 34-Figure 35). This knee is not present on the curve representing the reference material. The reason could be that the ageing has increased the sensitivity to variation in strain rate. The moment the tensile specimen begins to plasticise, it’s difficult to keep the strain rate absolutely constant because of the reaction in the material. A material sensitive to changes in strain rate could then exhibit this “knee” formation on the stress-strain curve.

51

As mentioned in chapter 3.2, the ordering effect leads to increased yield strength in nickel-based alloys while the ductility is deteriorated [29]. The results obtained in this study don’t reflect that behaviour. On the contrary, the yield strength is improved as well as the ductility after ageing. It was also reported [2] that the tensile strength and the ductility decreased with increasing ageing temperature due to ordering in Ni2Cr and Ni3Cr. That behaviour does not match the results in this study either. That the ductility is improved is also shown by the increased elongation and the larger area reduction of the tensile specimens. This result is in good agreement with the impact test results, even though the difference between material aged at 500°C and material aged at 550°C are not as pronounced as in the impact test. Hardness test was performed on samples on which no mechanical test was performed as well as the specimens from the LCF test. The specimens from the LCF test were included in the hardness testing in order to give some idea about how much the strain hardening, caused by the mechanical test, can affect the hardness of the material, but also the effect of the ageing temperature. The results from the macrohardness tests are found in Table 8. As expected, the strain hardening which occurs during the LCF test is strongly affecting the hardness in the unaged material, which increases with almost 40 % compared to the virgin bar. On the other hand, the hardness doesn’t seem to be affected appreciably by the ageing temperature in the strain hardened LCF specimens (see Figure 37). However, the untested specimens are affected by the ageing. As seen in Figure 36, the hardness increases with the ageing time. The hardness in Haynes 230 is generally attributed to the presence of carbides. An increase in hardness due to ageing could indicate on a precipitation of carbides during the heat treatment. This assumption agrees with the results from the microhardness test (see Figure 38) where the hardness near the grain boundaries increases due to ageing. It can be assumed that the precipitation of secondary carbides at grain boundaries occurring during ageing (for further discussion, see chapter 4.3.2) is at least partially responsible for this improvement. Since the microhardness test imprints were made in the matrix, the presence of primary carbides should not influence on the results. However, the microhardness test also showed that the hardness increases in the matrix in the middle of the grain as well, in an area where no precipitation of secondary carbides has been observed. Another interesting aspect is that the microhardness has its maximum after 1000 h of ageing both near grain boundaries and in the middle of the grain. The fraction of secondary carbides is not smaller after ageing at 550°C than at 500°C; it’s rather the opposite situation. These results give reason to believe that the hardness improvement is not only due to carbide precipitation. Whether the improvement after 1000 h of ageing depends on a hardening of the matrix, which is then reduced by continuous ageing, or if there are other factors responsible for this significant hardness variation is impossible to establish without further investigation. The fact that the material is harder after ageing for 1000 h than after 3000 h is in good agreement with the previous study on Haynes 230 carried out at SIT AB in Lincoln (see Figure 27 in chapter 3.2), although the ageing temperatures are lower in this study. When measuring the microhardness, the purpose was to avoid the visible carbides. Unfortunately it is impossible to locate the carbides under the surface,

52

which makes it probable that some of the performed imprints by accident have hit underlying carbides. These imprints are likely to have given a higher value of the hardness than expected. This is a likely reason to why the distribution of the results is larger in the microhardness test than in the macrohardness test. The higher test load applied the wider and deeper will the imprint be, which increases the risk of hitting underlying carbides.

4.3.2 Microstructural examination

The microstructural examination is based on an evaluation of images which are taken with LOM and SEM, chemical analyses measured by SEM and finally carbide size measurements. All images are found in Annex 8.6.1. An attempt was made to measure the carbide size in the aged and unaged Haynes 230 tensile test specimens. The result is given in Table 10. The measurements were made on micro photos taken with LOM at 500x magnification with Image-Pro Plus software. Because of limitations in the software and a resolution of the images which was inadequate, a large part of the carbides were excluded from the size measurement, which makes the result too unreliable to permit any conclusions about carbide size changes due to ageing. What can be established through a visual evaluation of the images together with the results in Table 10, without jumping to conclusions, is that the primary carbides are cellular to the shape and the average diameter is in the range of a few micrometers. The SEM was used to carry out the chemical analysis. Despite the fact that this SEM should be able to detect atoms with atomic number 5 and higher, including carbon, this is in reality not entirely true. The quantitative analysis of carbon is not done properly and a false value is returned. This error will of course affect the total composition measurement as well, to a small extent. More important is that the carbide composition determined by this SEM is unreliable since, in most cases, the returned value of the carbon content was equal to zero. This is of course wrong. Despite the chemical analysis, the exact carbon content in the matrix and the carbides is unknown. Another problem arose during the analysis of the secondary carbides. The depth of the impact of the electron beam in the material was greater than the depth of the secondary carbides. This implied that a part of the analysed material belonged to the matrix, which resulted in a composition that did not agree with the true composition of the secondary carbides. This is the reason why the conclusions from these results have been drawn with great cautious. The chemical analysis of the matrix is found in Table 11. No deviations from the references can be found. The exactness of the measurements isn’t good enough to make the decimals reliable. This explains the presence of titanium in the matrix of the specimen aged at 300 h at 550°C and the absence of cobalt in the untested specimen aged at 1000 h. As mentioned before, no attention should be paid to the 2.2 % of carbon either. No depletion of the matrix is discernable, not even after the longest ageing time or at any ageing temperature. Table 12 lists the results from the analysis of the primary carbides. The analysis is mainly performed on large carbides since the small ones had a tendency of being smaller than the depth of the electron beam impact, hence a part of the analysed volume belonged to the matrix. Just like the matrix, the composition of the primary carbides seems to be unaffected by the ageing at intermediate

53

temperatures. The dominating elements in the primary carbides are tungsten, nickel, chromium and molybdenum. Principally the primary carbides contain more tungsten and molybdenum than the matrix and less nickel, chromium, iron and manganese. As mentioned in chapter 2.1.2.3 the primary carbides are generally formed by a union between carbon and refractory elements [17], which in the case of Haynes 230 can be for example tungsten, molybdenum or chromium. The results from the analysis of secondary carbides are shown in Table 13. In general the composition is comparable to the composition of the matrix. However, the content of nickel, molybdenum and tungsten varies for the different specimens. Unfortunately it is difficult to find a tendency. Some specimens contain a larger part of nickel than the matrix, some contain less. The same pattern is repeated for the content of molybdenum and tungsten. Consequently it is impossible to decide the composition of secondary carbides based on these results. Figure 47 in Annex 8.6.1 is a survey image at low magnification of tensile tested, unaged Haynes 230. All the expected components are visible in Figure 48. The primary carbides are situated either at grain boundaries or intragranular. A selection of them is pointed out with white arrows. Twin boundaries, which frequently exist in Haynes 230, are pointed out with black arrows. Unfortunately, small areas are overetched, preferentially the grain boundaries which are extra vulnerable. An example is pointed out by the thick, white arrow. Some of the primary carbides break during tensile test (see Figure 49), creating a void around or between the carbide parts. It is desirable to avoid large void formations since they are causing fracture initiation. An edge fracture is visible in Figure 50. It looks like the crack propagation is prevented by the presence of the carbide. Overall there is no precipitation of secondary carbides in the unaged material. Figure 51 shows the microstructure of the specimen aged for 300 h at 500°C. The grain boundaries are more or less free from secondary carbides. Figure 52 shows that the majority of the grain boundaries are unaffected by the ageing, even though secondary carbides have precipitated at a small part of the boundaries. The twin boundaries on the other hand are not visually affected. In Figure 53 a slightly larger extent of the grain boundaries contain precipitated secondary carbides. Besides, there is a lot of void formation around the broken primary carbides. Figure 54 shows a crack at the fracture surface where the primary carbides are holding the crack together. The only visible ageing effect after 300 h at 500°C is a small extent of secondary carbides precipitated at grain boundaries. The effect of the ageing is slightly increased at 550°C. Figure 55 shows grains with unaffected boundaries while there are precipitated secondary carbides at the grain boundaries in Figure 56. Voids are formed around primary carbides just like for the reference. Figure 57 shows grain boundaries partially covered with secondary carbides. These three images illustrates a distribution of the primary carbides where all the large carbides are formed at grain boundaries or twin boundaries, while it is common to find the small carbides inside the grains. The effects of ageing at 500°C are larger after 1000 h than after 300, which is visible in Figure 58. The precipitation of secondary carbides is well distributed over the boundaries. The huge void formation around primary carbides visible at the image is probably created during preparation of the cut-up. Since the carbide probably was broken into pieces during tensile test, there is a risk that some of

54

the parts fell off during grinding and polishing of the cut-up. Figure 59 shows that there are grain boundaries free from precipitated secondary carbides in this specimen as well. The size of the primary carbides can influence on the fracture properties. Large primary carbides, as in Figure 60, run a greater risk of breaking when a load is applied than smaller carbides do. The larger the carbide is the larger will the void formation be. These internal microfractures can impair the rupture resistance. Figure 61 shows an edge crack restrained by primary carbide. Due to the hardness the carbide possesses, the propagation has to go around the carbide rather than through it. This will make the way longer for the crack and increase the energy absorbed during the crack propagation, which is positive. Figure 62 and Figure 63 shows the microstructure of the tensile test specimen aged at 1000 h at 550°C. There is a significant increase of secondary carbides compared to both the specimen with shorter ageing time and the specimen with lower ageing temperature. It is also visible that the void formation is present around medium carbides as well. Figure 64 shows a fracture that has propagated around the primary carbide. An increased amount of secondary carbides can, if they are properly formed, strengthen the grain boundaries and restrain grain boundary sliding [18]-[19]. The increased tensile strength of the specimen aged at 550°C for 1000 h can probably be at least partially attributed to the large amount of secondary carbides precipitated at grain boundaries. However the specimens aged at 500°C improved their tensile properties as well, so there is most likely yet another reason for the property change. Figure 65 and Figure 66 shows the microstructure of the virgin bar. Here it’s even more obvious that the grain boundaries are free from precipitations and that the large primary carbides are found at grain boundaries. It is also possible to look at the original shape of the primary carbides. The most frequent type is the cellular-shaped carbide which often is of small or medium size. The larger carbides are commonly of irregular shape (see Figure 67-Figure 68). Figure 69-Figure 70 represents the microstructure of untested specimen aged at 1000 h at 550°C. The first image shows no secondary carbides whatsoever but they are visible at the grain boundaries in the second image. An interesting observation was discovered when comparing this untested specimen with the tensile test specimen exposed to the same ageing conditions. The precipitation of secondary carbides has occurred to a much greater extent in the tensile test specimen than in the untested specimen. This observation indicates that the precipitation of secondary carbides can either be stimulated under an applied load or is simply easier to see after tensile testing. Figure 71 and Figure 72 shows that the larger primary carbides are located at the grain boundaries and that the intragranular carbides are much smaller to the size. Figure 73 and Figure 74 are from the untested specimen aged for 3000 h at 550 °C. They show that the precipitation of secondary carbides has increased due to the prolonged ageing time. There are no other visible effects of the ageing. In Figure 75 we can see that there are occasional large carbides at intragranular sites. The micro photos of the LCF tested specimens showed one unexpected phenomenon that distinguished these images from the rest. This phenomenon was discovered in specimens from all the different ageing conditions. Primary carbides, with what looks like holes in them, are shown in Figure 76 to Figure 81.

55

The study has not given any explanation to why these carbides are present in only the LCF tested specimens. Neither has it been established if the areas really are holes in the carbides or for example precipitations of a different phase. The samples where sent to Leicaster University in England for further investigations. The microstructural evaluation shows that the ageing at intermediate temperatures influences mainly on the precipitation of secondary carbides. It is shown that the precipitation was stimulated by prolonged ageing time but mainly by increased ageing temperature. These results correlate with the results from the study carried out at SIT in Lincoln [32], except that no precipitation at twin boundaries was visible at these intermediate ageing temperatures. The matrix composition measured in the virgin bar was utilised to calculate the Z-criterion. The obtained value was 1.62, which is smaller than 3.0, hence the LRO formation is possible according to the theory.

4.4 CONCLUSION

The ageing of Haynes 230 has a positive effect on the properties. The LCF properties proved to be unaffected, except for the maximum stress which is higher in the aged material. The same behaviour is found in the tensile test where the yield strength and tensile strength was improved due to ageing. The impact test and the elongation in the tensile test showed that the material became more ductile after the ageing. The macrohardness increased with the ageing temperature. The strain hardening of Haynes 230 highly improved the macrohardness and the ageing influence on the material decreased simultaneously. The microstructural examination neither confirmed nor refuted a growth of primary carbides. Neither could the composition of the secondary carbides be determined. The image analysis showed that the ageing has a positive effect on the precipitation of secondary carbides at grain boundaries, which normally leads to strengthened boundaries and restrained grain boundary sliding. The ageing temperature has a larger effect on the microstructure than the ageing time. The Z-criterion indicates that the existence of LRO in Haynes 230 is probable. The fact that the material has proved to shrink during ageing is another indication. Haynes 230 has many similarities to Hastelloy in which the LRO has been observed [2]. These indications are enough convincing to justify further research in this matter.

56

5 AGEING INFLUENCE ON MECHANICAL PROPERTIES AND MICROSTRUCTURE OF CMSX-4 AT INTERMEDIATE TEMPERATURES (500-550°C)

5.1 SPECIMEN PREPARATION AND EXPERIMENTAL PROCEDURE

5.1.1 Tensile test

Ten tensile test specimens were cut out from 5 different bars, two specimens from each bar. Each pair had the same single crystal direction. The dimensions of the bars followed the drawing 7000 1782-3 which is found in Annex 8.3.3. One specimen from each pair was aged at 500°C during 1150 h. Data for each specimen showing direction and angles for the single crystal bar, together with ageing time and temperature, are listed in Table 14.

Specimen marking

SX direction

Theta Rho Ageing temp [°C]

Ageing time [h]

114a Off angle 25.9 4.5 500 1150 114b Off angle 25.9 4.5 - - 122a Off angle 19.8 11.3 500 1150 122b Off angle 19.8 11.3 - - 124a <011> 42 0.2 500 1150 124b <011> 42 0.2 - - 126a <011> 41.9 7.6 500 1150 126b <011> 41.9 7.6 - - 143a <111> 53.8 43.8 500 1150 143b <111> 53.8 43.8 - -

Table 14: Specimen data for CMSX-4.

The tensile tests were performed in the tensile test machine Schenck-Trebel RSA 100 on all specimens in agreement with EN 10 002-1. The strain was measured by an extensometer up to 2.0 %. The diameter of the specimens was measured before and after the testing. Tensile strength was obtained from the software calculations. The elongation of the bar and the reduction of the cross-section area were calculated together with the yield strength at 0.1 % and 0.2 % plastic strain deformation. All details are given in Annex 8.4.2.

5.1.2 Microscopic examination

After the tensile test, the specimens were photographed in order to document the fracture surfaces and determine the angle between the slip planes and the test load axis. How to obtain the angle is illustrated in Figure 40.

57

Afterwards, cut-ups containing the tensile test specimens were prepared, exposing the cross-section along the test load axis. The same procedure was employed as in the case of Haynes 230 (see chapter 4.1.4) except for the etching. For this purpose, a mixture containing MoO3, HCl, HNO3 and H2O was used. The CMSX-4 cut-ups were etched for approximately 25 s. The cut-ups were examined in both optical microscope and SEM. The focus was set on any differences between new and aged material, for example on the appearance of active slip planes and the distance between them.

Figure 40: Angle of slip plane, specimen 126a.

5.1.3 Hardness test

Because of the structure of CMSX-4, it wasn’t suitable to perform the hardness test on the cut-ups. The reason is that it was impossible to confirm that the SX direction in the hardness test loading direction was the same for both specimens in a pair. Rotation of the cylindrical specimen would alter the SX direction perpendicular to the ground and polished cross-section. Since CMSX-4 isn’t isotropic, a comparison between the two specimens in a pair would in that case be pointless. The macrohardness test was therefore performed on the cross-section of the piece that was cut off during the preparation of the cut-up. This was the only way to guarantee the same SX direction for both specimens in a pair and thus be able to compare the results of new and aged material. The equipment used for macrohardness test for CMSX-4 is the same used in the previous case of Haynes 230 (see chapter 4.1.5). The applied load was 10 kg and the hardness was measured in Vickers scale (HV10). The distance between the imprints followed the same rules used in the macrohardness testing of Haynes 230. See chapter 4.1.5 for more details.

58

5.2 MECHANICAL TEST RESULTS

5.2.1 Tensile test

Nine of the tests were successful and one test (specimen 114a) was stopped before fracture. The results from the tensile tests are shown in Table 15. The difference between new and aged material is calculated for each pair and property.

Aged at 1150h at

500°C

New material

(ref)

Rm [MPa]

Rp0.1 [MPa]

Rp0.2 [MPa]

A5 [%]

Z [%]

114b 833 810 817 6.8 30.2 114a 756 742 747 - - Difference [%] -9.2 -8.4 -8.6 - -

122b 873 872 871 11.7 28.3 122a 796 783 783 13.2 31.3 Difference [%] -8.9 -10.2 -10.1 12.8 10.6

124b 1000 1000 1000 3.5 - 124a 885 869 869 15.0 33.0 Difference [%] -11.5 -13.1 -13.1 329 -

126b 924 924 924 12.2 29.0 126a 867 858 858 10.4 34.4 Difference [%] -6.1 -7.1 -7.1 -14.8 18.6

143b 1465 1146 1156 9.6 6.4 143a 1490 1154 1155 - 10.4 Difference [%] 1.8 0.7 -0.1 - 62.5

Table 15: Tensile test result for CMSX-4.

The stress-strain curves plotted up to 2 % strain, or until failure if it occurred at lower strain, are presented in Figure 41 to Figure 45.

59

0

2 0 0

4 0 0

6 0 0

8 0 0

1 0 0 0

0 0 , 2 0 , 4 0 , 6 0 , 8 1 1 , 2

S t r a i n ( % )

S t r e s s( M P a )

1 1 4 a - a g e d 1 1 5 0 h 1 1 4 b

Figure 41: Stress-strain curves from tensile test, specimen pair 114.

0

2 0 0

4 0 0

6 0 0

8 0 0

1 0 0 0

0 0 , 1 0 , 2 0 , 3 0 , 4 0 , 5 0 , 6

S t r a i n ( % )


1 2 2 a - a g e d 1 1 5 0 h 1 2 2 b


0

2 0 0

4 0 0

6 0 0

8 0 0

1 0 0 0

1 2 0 0

0 0 , 1 0 , 2 0 , 3 0 , 4 0 , 5

S t r a i n ( % )


1 2 4 a - a g e d 1 1 5 0 h 1 2 4 b


60

0

2 0 0

4 0 0

6 0 0

8 0 0

1 0 0 0

0 0 , 1 0 , 2 0 , 3 0 , 4 0 , 5

S t r a i n ( % )


1 2 6 a - a g e d 1 1 5 0 h 1 2 6 b


0

2 0 0

4 0 0

6 0 0

8 0 0

1 0 0 0

1 2 0 0

1 4 0 0

0 0 , 5 1 1 , 5 2 2 , 5

S t r a i n ( % )


1 4 3 a - a g e d 1 1 5 0 h 1 4 3 b


61

5.2.2 Angle of slip plane

The angle of the slip planes is presented in Table 16.

Specimen marking

SX direction

Theta [°] Angle of slip plane

114a Off angle 25.9 No fracture 114b Off angle 25.9 42.5 122a Off angle 19.8 45 122b Off angle 19.8 42 124a <011> 42 45.5 124b <011> 42 44.5 126a <011> 41.9 50.5 126b <011> 41.9 35 143a <111> 53.8 - 143b <111> 53.8 -

Table 16: Angle of slip planes after tensile test.

The determination of the angle of the slip planes of pair 143 and 114a was impossible, because of the lack of visible marks on the specimen surface.

5.2.3 Hardness test

The results from the macrohardness testing are shown in a chart in Figure 46 together with the difference between new and aged material.

M a c r o h a r d n e s s

- 6 . 8 % - 2 . 7 % - 0 . 6 % 3 . 2 %

6 . 0 %

3 0 0

3 5 0

4 0 0

4 5 0

5 0 0

5 5 0

6 0 0

1 1 4 1 2 2 1 2 4 1 2 6 1 4 3

S p e c i m e n N o

H a r d n e s s[ H V 1 0 ] n e w

a g e d

Figure 46: Results from macrohardness test for CMSX-4.

62

All data from the macrohardness tests are visualised in Table 17.

Specimen Pair No

Status Vickers hardness[HV10]

Standard deviation


New 424 8.87 114

Aged 395 4.09 -6.8 New 444 10.9

122 Aged 432 17.5 -2.7 New 434 3.83

124 Aged 431 8.73 -0.6 New 423 5.16

126 Aged 437 9.40 3.2 New 516 8.57

143 Aged 547 17.4 6.0

Table 17: Data from macrohardness testing for CMSX-4.

5.3 DISCUSSION

5.3.1 Mechanical tests

Drawing conclusions from tensile tests performed on single crystal material is always a bit hazardous. The deformation is concentrated to a few active slip planes and whether these slip planes occur within the extensometer range or not have a crucial effect on the deformation curves, i.e. the stress-strain curves (see Figure 41-Figure 45). If the crack initiation takes place outside the extensometer range the deformation is concentrated to that area and the strain measured by the extensometer will be much smaller than the accurate value, i.e. the curve is not appropriate. This is the reason why tensile test is not an appropriate method to determine neither the elastic modulus nor the tensile strength at failure, and these results are therefore not discussed in this study. The value of the elongation at break is not taken from the curves. Instead the specimen bars were measured after finished test. A distinct trend in comparison between new and aged material is the reduction in yield strength and tensile strength (see Table 15). After being aged during 1150 h, the decrease of Rp0.1 and Rp0.2, which followed each other very well, was found to be between 7 and 13 %. The reduction in tensile strength as well was found to bee in good correlation with these results. The only exception among the specimens is the specimen pair in direction <111>, where the change due to ageing was very small or negligible. The deformation of the specimens developed during tensile loading depends on the direction of the single crystal. The specimens in direction <111> (pair 143) developed no neck during loading. On the contrary, the active slip planes were uniformly distributed over the whole length of the bar, resulting in a slightly elliptic cross-section of the bar. The tensile tests with the specimens in direction

63

<011> or off angle resulted in elliptic cross-sections as well, but these were far more flattened. Furthermore, the slip planes were concentrated to a small part of the bar, resulting in a neck formation in this area. Common for all the specimens was that the larger diameter of the elliptic cross-section was not affected by the tensile test. The effect of the ageing phenomenon on the deformation is palpable. The overall result is that the aged specimens show a larger elongation and a greater reduction in cross-section than the specimens which were not aged. One exception is found for specimen pair 126 where the ageing resulted in a smaller elongation of the specimen. The magnitude of the deformation values varies a lot. One reason could be that the measuring was made manually with help of a slide calliper, a fact that can have a negative effect on the accuracy of the measurements. Another possible reason is the lack of statistical basic data. An average from a number of tensile tests carried out under the same ageing and testing conditions would probably give a result more faithful to the true value. It’s difficult to find a trend in the hardness results (see chapter 5.2.3) due to ageing. Pair 114 (theta = 25.9°) shows a decrease in hardness with 6.8 % while pair 143 (SX direction <111>) shows an increase of 6.0 % due to ageing. The other results are well distributed between those values. However, noteworthy is that the hardness seems to depend on the SX direction. The pairs off direction or in the direction <011> have an average between 400 and 450 HV10. Pair 143 on the other hand has an average closer to 550 HV10. The slip plane angle was impossible to determine for three of the specimens due to difficulties to see the marks from the slip planes on the surface of the specimens. The results from the remaining three pairs reveal no distinct trend (see Table 16). The slip plane angle is 3 degrees larger after ageing for pair 122 (theta = 19.8°). The difference for pair 124 (theta = 42°) is only 1 degree. In pair 126 (theta = 41.9°) the difference is approximately 15 degrees. Specimens 124 and 126 have more or less the same SX direction. Therefore a similar slip plane angle for these specimens was expected. However, specimen 126b measured a much smaller slip plane angle compared to the other ones. Some of the results for specimen 126b are quite unexpected, such as the small elongation in tensile test and the angle of the slip plane. Due to the deviation from the expected results the conclusions regarding the performance of specimen pair 126 should be drawn with cautiousness. The high temperature properties of CMSX-4 are thoroughly investigated in a great number of studies. However, there is no documentation of the behaviour at intermediate temperatures, which makes it difficult to estimate the accuracy of the results obtained in this study.

5.3.2 Microstructural Examination

The microstructural examination is based on images which are taken with LOM and SEM. All images are found in Annex 8.6.2. No microstructural examination is performed on pair 114. Since the tensile test of specimen 114a was stopped prematurely, hardly any deformation had time to occur. A comparison of the pair would consequently be pointless since the conditions weren’t the same for both specimens. Figure 82 shows the microstructure of CMSX-4. The image is taken from specimen 122a, which was aged before the tensile test. The dark zones are γ´

64

particles. They are embedded in the matrix which appears as a white check pattern on the SEM photos. The image is taken at 4000x magnitude. The slip planes activated during the tensile test are recognised as light stripes in the material. Figure 83 and Figure 84 from specimen 122a shows that there are two different types of distribution of active slip planes. The first case is when single slip planes are activated and they are distributed with a smaller or larger distance to the next slip plane. The second case is when one slip plane is activated and the adjacent slip planes follow after. This behaviour results in a band of activated slip planes. An enlargement of the slip planes (see Figure 85 and Figure 86) reveals that not only the matrix is deformed during the tensile test but also the gamma primes. The images show slip planes where the gamma primes have been strongly prolonged. Figure 87 shows an enlarged image of a band of activated slip planes. The aged specimen 122a had a rather large number of activated slip planes compared to the unaged specimen in the pair, 122b. The only slip planes found for that specimen are shown on a LOM image in Figure 88. Figure 89 shows the same area reproduced by the SEM. The lower surface is the edge of the specimen and the left surface is the fracture surface in Figure 89. An interesting remark is that the deformation in the material preferably occurs as atom planes sliding against each other and not as crack propagation. Figure 90 shows crack initiations on the edge of the specimens, which have not led to fracture. In Figure 91 an enlargement of single slip plane with deformed gamma prime is visible and in Figure 92 a narrow band of slip planes shows how the material is deformed under traction. The difference between 122a and b is that the amount of activated slip planes is larger for the aged specimen. It also contains more bands of slip planes while the unaged specimen mostly contains single slip planes separated from each others. No favourable distance between the slip planes was discerned. On the contrary, the distribution of the slip planes seems arbitrary. Specimen 124a contained only a few activated slip planes, all widely separated from each other. Figure 93 shows some feeble marks from single slip planes, but also a few unwanted inclusions of unknown kind. In Figure 94 and Figure 95 some single slip planes are visible. The enlargements in Figure 96 and Figure 97 does not show any deformed gamma primes that were visible in pair 122. The microstructure of specimen 124b was more or less the same. Some feeble marks of slip planes are visible in Figure 98, and the enlargement in Figure 99 doesn’t show any distinct deformation. The microstructure is slightly more affected around a crack tip (see Figure 100), but there is no sharp contrast between deformed and unaffected regions. There are no large differences between new and aged material in pair 124. They both exhibit only a few separate activate slip planes and the deformation doesn’t seem to affect the form of the gamma primes appreciably. There doesn’t seem to be any tendency concerning the distance between the slip planes in any of the specimens. The deformation is creating both separated slip planes and bands of slip planes in the aged specimen 126a (see Figure 101-Figure 103). The enlargements in Figure 104 and Figure 105 show that it’s rather the matrix than the gamma primes that is affected and deformed due to the glide. It is also shown in Figure 106 and Figure 107 that slip planes are possible crack initiators.

65

The images of unaged 126b (see Figure 108 and Figure 109) show a similar behaviour. Single slip planes are found next to deformed bands. The activated slip planes are well distributed over the specimen. Looking at Figure 110, the gamma primes seems unaffected by the deformation. Both specimens in pair 126 contain a lot of activated slip planes, both separated and in bands. The deformation seems to be concentrated to the matrix and there are no visible differences between new and aged material. The microstructure in pair 143 is, because of the SX direction, not at all like the microstructure of the other pairs. Since the loading axis is in the <111> direction, there are no natural slip planes where the gliding can occur. That’s why the behaviour seen in Figure 111 occurs. The crack propagation in the aged specimen 143a suddenly changes direction and continues along another path. An enlargement of the crack tip (see Figure 112) shows that the microstructure around the tip is more or less homogeneous, which makes it impossible to divine the direction of further propagation. The few visible slip planes (see Figure 113) resulted either in an internal fracture along the slip plane (Figure 114) or in a deformation of the matrix (Figure 115) comparable to the behaviour of pair 126 (see Figure 104). The unaged specimen 143b shows the same behaviour as the previously discussed specimen. Cracks, which changed propagation direction during the test (see Figure 116), were found in this specimen as well and the microstructure around the crack tip seemed homogeneous and without defaults (Figure 117). The traction resulted in internal fracture where the matrix was the weakest link in the chain and broke down (see Figure 118 and Figure 119), but also in deformation where the glide of atom planes deformed the matrix (see Figure 120 and Figure 121) as for the aged specimen in the pair. Once again it’s difficult to find any differences between new and aged material. The deformation expresses itself in the same ways in both specimens and the crack propagation is hazardous and unpredictable. The composition of the gamma matrix after precipitation of gamma prime has to be known to determine the Z-criterion. No chemical analysis of the specific material used in this study has been performed. Rothová et al. [33] have listed the chemical composition of γ and γ´ in CMSX-4 with 60 vol% γ´. According to them, the formula of gamma matrix after precipitation of gamma prime is Ni50Cr22Co20Re8. This composition renders in a Z-criterion equal to 2.27, which would imply that the LRO formation in CMSX-4 is probable.

5.4 CONCLUSION

No obvious differences in microstructure between new and aged material have been found that can explain the change in mechanical properties. A conceivable cause is the amount of gamma prime before and after the heat treatment. It is know that a too large fraction of gamma prime will deteriorate for example the ductility and the thermo mechanical fatigue properties. Simultaneously a too small fraction of gamma prime can result in a reduction in high temperature strength and creep resistance. Since no measurement of gamma prime fraction has been carried out within this study, no certain conclusions about whether it has changed or not can be drawn.

66

It is clear that ageing at 500°C influences on the mechanical properties. Generally the ageing results in a deteriorated tensile strength and an improved ductility. However, the impact the ageing has on the material is depending on the SX direction. The specimen pair in <111> direction had originally significantly higher tensile strength compared to the other specimens, and it was hardly affected by the ageing. The remaining specimens were all influenced by the ageing. This study has not brought out any proof of the existence of SRO/LRO in CMSX-4, even if the obtained Z-criterion indicates that the LRO formation is possible. It has been proved that the ageing has influence on the properties, but whether the material shrinks as a result of the ageing or not hasn’t yet been examined. However, nothing that tells against that the ordering phenomenon could occur in CMSX-4 has been found either.

67

6 FUTURE WORK

There is no statistical basis for the mechanical tests performed on CMSX-4. Testing of tensile properties and hardness with an enlarged number of specimens would give a more reliable result. In other studies the variation of the carbide size in Haynes 230 has been correlated to property changes. A successful measuring of the carbide size would facilitate the investigation of microstructural changes and its impact on the mechanical properties. Hopefully it would result in a better understanding of optimal heat treatment temperatures and times for Haynes 230. An increased volume fraction of gamma prime in CMSX-4 is a possible explanation to the property changes after ageing. By measuring the volume fraction the understanding of the role gamma prime plays in CMSX-4 would be ameliorated. Until today, no measurements of shrinkage in CMSX-4 have been performed. A negative result from such test could probably dismiss the ordering phenomenon theory as an underlying cause to the property changes in the single crystal material. The presence of LRO in the materials can be proved by finding Ni2Cr superlattice reflections in electron diffraction and an accompanying decrease in electrical resistivity.

68

7 REFERENCES

[1]. Moverare, J., (2005). Shrinkage and ageing of Ni-base alloys at intermediate temperatures (400-600°C). Finspång: Siemens Industrial Turbomachinery AB. Internal technical report 1CS51497.

[2]. Marucco, A., Nath, B., (1988). Effects of ordering on the properties of Ni-Cr alloys. J. Mat. Sci., 23, 2107—2114.

[3]. Granacher, J., Pfenning, A., (1994). Optimierung und Verifizierung von Kriechgleichungen für Hochtemperaturwerkstoffe. Darmstadt: Institut für Werkstoffkunde der Technischen Hockschule Darmstadt Technical report 523.

[4]. SGT-800 presentation from SIT AB.

[5]. General Electric Reference Document: Gas Turbine Design Philosophy. GER-3434D.

[6]. http://www.grc.nasa.gov/WWW/K-12/airplane/brayton.html

[7]. http://en.wikipedia.org/wiki/Gas_turbine

[8]. Annerfeldt, M., et al, (2004). Result from the thermo-crystal test performed in B520 Helsingborg Blades and Vanes temperature and SAS pressure. Finspång: Siemens Industrial Turbomachinery AB. Internal technical report RT T10C 50/03.

[9]. General Electric Reference Document: Advanced Gas Turbine Materials and Coatings. GER-3569G.

[10]. Shukin, S., Björkman, M. SGT-800 Vertion A+. Turbine Blade 1 Design, 3D Thermal & MI analysis. Finspång: Siemens Industrial Turbomachinery AB. Internal technical report RT GRC 249/05.

[11]. Orakzai, K., (2004). Effect of Carbon Content on the Properties of CMSX-4. Finspång: Siemens Industrial Turbomachinery AB. Internal technical report 1CS41538.

[12]. Chapman, L.A., (2004). Application of high temperature DSC technique to nickel based superalloys. J. Mat. Sci., 39, (24) 7229—7236.

[13]. Haynes 230 alloy Product Brochure H-3000H.

[14]. Hogmark, S., Jacobson, S., Kassman-Rudolphi, Å., (1998). Elektronmikroskopi i praktik och teori. Uppsala: Uppsala universitet, Ångströmslaboratoriet.

[15]. http://www.rpi.edu/dept/materials/COURSES/NANO/shaw/Page5.html

[16]. Goodhew, P.J., Humphreys, F.J., (1988). Electron Microscopy and Analysis. 2nd edition. London: Taylor & Francis.

[17]. Sims, C.T., Stoloff, N.S., Hagel, W.C., (1987). Superalloys II. New York: Wiley-Interscience.

[18]. Donachie, M.J., Donachie, S.J., (2002). Superalloys. A Technical Guide. 2nd edition. Materials Park: ASM International.

[19]. Davis, J.R., (1997). Heat Resistant Materials. Materials Park: ASM International.

[20]. http://www.msm.cam.ac.uk/phase-trans/2003/Superalloys/superalloys.html

http://www.grc.nasa.gov/WWW/K-12/airplane/brayton.html

http://en.wikipedia.org/wiki/Gas_turbine

http://www.rpi.edu/dept/materials/COURSES/NANO/shaw/Page5.html

http://www.msm.cam.ac.uk/phase-trans/2003/Superalloys/superalloys.html

69

[21]. http://www.msm.cam.ac.uk/phase-trans/2003/Superalloys/rafting.html

[22]. Material certificate from Cannon-Muskegon, Heat No: V 8508 and V 8677.

[23]. Johnsson, A., (2001). A LTA/LTE study on the thermal characteristics of Haynes alloy 230. Finspång: Siemens Industrial Turbomachinery AB. Internal technical report RT TRM 54/98.

[24]. Marucco, (1989). Materials Architecture. 10th Risø International Symposium on Metallurgy and Materials Science

[25]. Marucco, A., Nath, B., (1991). Effects of γ´ volume fraction on ordering behaviour and rupture properties of Alloy 80A. Proc. 6th International Conference on Mechanical Behaviour of Materials (ICM6), Kyoto, Japan.

[26]. Metcalfe, E., Nath, B., Wickens, A., (1984). Some Effects of the Ordering Transformation in Nimonic 80A on Stress Relaxation Behaviour. Mat. Sci. Eng., 67, (2), 157—162.

[27]. Reppich, B., (1994). Negatives Kriechen und Mikrogefüge langzeitexponierter Gasturbinenwerkstoffe. Z. Metallkd., 85, (1).

[28]. Marucco, A., (1995). Phase transformations during long-term ageing of Ni-Fe-Cr alloys in the temperature range 450-600 °C. Mat. Sci. Eng., A194, (2), 225—233.

[29]. Rtishchev, V.V., (1994). Structure transformations and property changes of Ni-base superalloys on ageing. Mat. Adv. Power Eng., (1), 889—898.

[30]. Pettinari, F., Prem, M., Krexner, G., Caron, P., Coujou, A., Kirchner, H.O.K., Clément, N., (2001). Local order in industrial and model γ phases of superalloys. Acta Mat., 49, 2549—2556.

[31]. http://www.cmse.ed.ac.uk/AdvMat45/SuperMicro.pdf

[32]. Marchant, G., (2005). Microstructural degradation of combustor alloy Haynes 230. Lincoln: Siemens Industrial Turbomachinery AB. Internal technical report 05-LAB-211

[33]. Rothová, V., Stloukal, I., Čermák, J., (2000). Permeation of hydrogen in Ni-based superalloy CMSX-4. Acta Mat., 48, 827—833.

http://www.msm.cam.ac.uk/phase-trans/2003/Superalloys/rafting.html

http://www.cmse.ed.ac.uk/AdvMat45/SuperMicro.pdf

70

8 ANNEXES

8.1 MATERIAL DATASHEET FOR HAYNES 230 COATING

71

8.2 TEST CERTIFICATE, HAYNES 230, HEAT NO 1830587801

72

8.3 DRAWINGS

8.3.1 Low cycle fatigue test specimen

73

8.3.2 Impact test specimen

74

8.3.3 Tensile test specimen

75

8.4 CALCULATIONS

8.4.1 LCF test

Ni and Nf – Number of cycles to crack initiation and failure Ni is defined as 5 % load drop from the stabilised loop. A linear interpolation of the loop between 30 % and 70 % of the total amount of cycles was compared to the actual stress in each point. Ni was obtained when the drop exceeded 5 %. The same method was used to calculate Nf, which is defined as 50 % load drop.

σmax (N = 1) – Maximum stress at first cycle The maximum stress at the first cycle was read off during the test.

σmax (Nf / 2) – Maximum stress at midlife The maximum stress at midlife was obtained from the graph where maximum stress was plotted against the number of cycles.

Δεpl (Nf / 2) – Plastic strain range at midlife Δεpl was read off the graph where plastic strain was plotted against the number of cycles.

8.4.2 Tensile test

A5 – Elongation of waist of specimen

10050

0 ⋅−

=LLL

A Equation 4: Elongation of waist of test specimen.

L0 is the initial length between two fixed points at the waist of the specimen and L is the length between the two points after tensile test. Z – Reduction of cross-section area of specimen

100

4

4420

220

⋅−

=D

DD

Zπ

ππ

Equation 5: Reduction of cross-section area of specimen.

D is the diameter of the waist after tensile test and D0 is the initial diameter. Rp0.1 – Yield strength at 0.1 % plastic strain deformation

A linear curve with the same slope as the stress-strain curve and with an offset of 0.1 % strain was plotted. The point where the stress-strain curve and the linear curve intersected was defined as the yield point at 0.1 % plastic strain deformation. The stress read off in the graph at this point is defined as the yield strength.

76

8.5 RESULTS

8.5.1 Tensile test, Haynes 230

Ref 300 h 1000 h 2500 h 500°C 550°C 500°C 550°C 500°C 550°C

Rm [MPa] 862 871 869 871 873 876 873 Rp0.1 [MPa] 366 400 400 398 401 403 415 Rp0.2 [MPa] 379 395 392 394 397 400 400

A5 [%] 47.8 50.7 53.7 50.6 50.7 52.5 51.8 Z [%] 48.0 55.7 57.6 54.3 50.1 55.8 55.8

Table 18: Result from tensile test, Haynes 230.

77

8.6 MICROSCOPY PHOTOS

8.6.1 Haynes 230

Figure 47: Tensile test specimen ref – survey image taken with LOM. Marker size: 100 μm.

Figure 48: Tensile test specimen ref – SEM photo of twin boundaries (black arrows) and primary carbides (white arrows). The thick white arrow shows an overetched area. Marker size: 50 μm.

78

Figure 49: Tensile test specimen ref – SEM photo of void formation around broken primary carbides (black arrows). Marker size: 50 μm.

Figure 50: Tensile test specimen ref – SEM photo of edge fracture stopped by primary carbide (black arrow). Marker size: 20 μm.

Figure 51: Tensile test specimen, aged 300 h at 500°C – SEM photo of grain boundaries free from secondary carbides. Marker size: 50 μm.

79

Figure 52: Tensile test specimen, aged 300 h at 500°C – SEM photo of twinning and grain boundaries partially free from secondary carbides. Marker size: 50 μm.

Figure 53: Tensile test specimen, aged 300 h at 500°C – SEM photo of grain boundaries with secondary carbides (black arrow) and void formation around broken primary carbides (white arrows). Marker size: 50 μm.

Figure 54: Tensile test specimen, aged 300 h at 500°C – SEM photo of fracture kept together by primary carbides (black arrows). Marker size: 50 μm.

80

Figure 55: Tensile test specimen, aged 300 h at 550°C – SEM photo of grain boundaries without secondary carbides. Marker size: 50 μm.

Figure 56: Tensile test specimen, aged 300 h at 550°C – SEM photo of secondary carbides precipitated at grain boundaries (black arrow). Marker size: 50 μm.

Figure 57: Tensile test specimen, aged 300 h at 550°C – SEM photo of void formation around broken primary carbides (black arrows) and precipitated secondary carbides (white arrows). Marker size: 50 μm.

81

Figure 58: Tensile test specimen, aged 1000 h at 500°C – SEM photo of large void formations around primary carbides (black arrows) and secondary carbides precipitated at grain boundaries (white arrows). Marker size: 50 μm.

Figure 59: Tensile test specimen, aged 1000 h at 500°C – SEM photo of grain boundaries with (black arrow) and without (white arrow) secondary carbides. Marker size: 50 μm.

Figure 60: Tensile test specimen, aged 1000 h at 500°C – SEM photo of large void formation around broken primary carbides (black arrows). Marker size: 50 μm.

82

Figure 61: Tensile test specimen, aged 1000 h at 500°C – SEM photo of primary carbide preventing crack propagation. Marker size: 20 μm.

Figure 62: Tensile test specimen, aged 1000 h at 550°C – SEM photo of an advanced stage of precipitation of secondary carbides at grain boundaries. Marker size: 100 μm.

Figure 63: Tensile test specimen, aged 1000 h at 550°C – SEM photo of void formation around primary carbides at grain boundaries. Marker size: 50 μm.

83

Figure 64: Tensile test specimen, aged 1000 h at 550°C – SEM photo of crack propagation without the presence of primary carbides (black arrow). Marker size: 50 μm.

Figure 65: Untested specimen, ref – SEM photo of grain boundaries without secondary carbides and large primary carbides at grain boundaries. Marker size: 50 μm.

Figure 66: Untested specimen, ref – SEM photo of small, intragranular primary carbides (black arrows) and twinning (white arrows). Marker size: 50 μm.

84

Figure 67: Untested specimen, ref – SEM photo of large, primary carbides at grain boundaries. Marker size: 50 μm.

Figure 68: Untested specimen, ref – SEM photo of distribution of primary carbides. Marker size: 50 μm.

Figure 69: Untested specimen, aged 1000 h at 550°C – SEM photo of twinning and grain boundaries without secondary carbides. Marker size: 20 μm.

85

Figure 70: Untested specimen, aged 1000 h at 550°C – SEM photo of precipitation of secondary carbides at grain boundaries (black arrows). Marker size: 50 μm.

Figure 71: Untested specimen, aged 1000 h at 550°C – SEM photo of distribution of primary carbides. Marker size: 50 μm.

Figure 72: Untested specimen, aged 1000 h at 550°C – SEM photo of small, intragranular primary carbides. Marker size: 50 μm.

86

Figure 73: Untested specimen, aged 3000 h at 550°C – SEM photo of small, intragranular primary carbides and twinning. Marker size: 50 μm.

Figure 74: Untested specimen, aged 3000 h at 550°C – SEM photo of precipitation of secondary carbides at grain boundaries. Marker size: 50 μm.

Figure 75: Untested specimen, aged 3000 h at 550°C – SEM photo of larger intragranular primary carbides. Marker size: 50 μm.

87

Figure 76: LCF tested specimen, ref – SEM photo of primary carbide with holes? Marker size: 10 μm.

Figure 77: LCF tested specimen, ref – SEM photo of primary carbides with holes? Marker size: 20 μm.

Figure 78: LCF tested specimen, aged 1000 h at 500°C – SEM photo of primary carbides with holes? Marker size: 10 μm.

88




89

8.6.2 CMSX-4

Figure 82: 122a – SEM photo of CMSX-4 structure. Marker size: 10 μm.

Figure 83: 122a – SEM photo of slip plane. Marker size: 100 μm.

Figure 84: 122a – SEM photo of slip plane. Marker size: 100 μm.

Single slip plane

Band of activated slip planes.

90

Figure 85: 122a – SEM photo of band of activated slip planes with deformed gamma prime. Marker size: 10 μm.

Figure 86: 122a – SEM photo of slip plane with deformed gamma prime. Marker size: 10 μm.

Figure 87: 122a – SEM photo of band of activated slip planes with deformed gamma prime. Marker size: 10 μm.

91

Figure 88: 122b – LOM photo of the specimen edge near fracture surface. Marker size: 20 μm.

Figure 89: 122b – SEM photo of separate slip planes at fracture surface. Marker size: 100 μm.

Figure 90: 122b – SEM photo of crack initiations. Marker size: 100 μm.

92

Figure 91: 122b – SEM photo of slip plane. Marker size: 10 μm.

Figure 92: 122b – SEM photo of band of activated slip planes with deformed gamma prime. Marker size: 10 μm.

Figure 93: 124a – SEM photo of microstructure with enclosures. Marker size: 100 μm.

93

Figure 94: 124a – SEM photo of single slip plane. Marker size: 50 μm.

Figure 95: 124a – SEM photo of separate slip planes. Marker size: 50 μm.


94


Figure 98: 124b – SEM photo of feeble marks of slip planes. Marker size: 100 μm.

Figure 99: 124b – SEM photo of single slip plane. Marker size: 10 μm.

95

Figure 100: 124b – SEM photo of structure around a crack. Marker size: 10 μm.

Figure 101: 126a – LOM photo of separate slip planes. Marker size: 50 μm.

Figure 102: 126a – LOM photo of separate and bands of activated slip planes. Marker size: 100 μm.

96

Figure 103: 126a – LOM photo of band of activated slip planes. Marker size: 50 μm.



97

Figure 106: 126a – SEM photo of fracture caused by gliding. Marker size: 50 μm.

Figure 107: 126a – Enlargement of deformation around crack tip in Figure 106. Marker size: 10 μm.

Figure 108: 126b – LOM photo of separate slip planes. Marker size: 100 μm.

98

Figure 109: 126b – LOM photo of band of activated slip planes. Marker size: 50 μm.

Figure 110: 126b – SEM photo of slip plane. Marker size: 20 μm.

Figure 111: 143a – SEM photo of crack which changed direction during propagation. Marker size: 50 μm.

99

Figure 112: 143a – SEM photo of structure around crack tip. Marker size: 10 μm.


Figure 114: 143a – SEM photo of internal fracture by slip plane. Marker size: 10 μm.

100


Figure 116: 143b – SEM photo of crack which changed direction during propagation. Marker size: 50 μm.

Figure 117: 143b – SEM photo of structure around crack tip. Marker size: 10 μm.

101

Figure 118: 143b – SEM photo of internal fracture by slip plane. Marker size: 10 μm.

Figure 119: 143b – SEM photo of internal fracture. Marker size: 10 μm.

Figure 120: 143b – SEM photo of deformation due to activated slip planes. Marker size: 10 μm.

102

Figure 121: 143b – SEM photo of deformation due to activated slip planes. Marker size: 10 μm.

Ageing influence on nickel-based superalloys at ...1022216/FULLTEXT01.pdf · these superalloys have...

Documents

Transcript of Ageing influence on nickel-based superalloys at ...1022216/FULLTEXT01.pdf · these superalloys have...