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Understanding NACE Understanding NACE MR0175-Carbon Steel Written Exam GlRdi General Reading. Reading 7 Fion Zhang/ Charlie Chong 3 rd Nov 2017 Fion Zhang/ Charlie Chong

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Oil And Gas Production Industry Fion Zhang/ Charlie Chong

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Oil And Gas Production Industry Fion Zhang/ Charlie Chong

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Oil And Gas Production Industry Fion Zhang/ Charlie Chong

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Fion Zhang/ Charlie Chong

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过五关斩六将 Fion Zhang/ Charlie Chong

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过五关斩六将 Fion Zhang/ Charlie Chong

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NACE MR0175 Carbon Steel NACE MR0175-Carbon Steel Written Exam Written Exam NACE-MR0175-Carbon Steel -001 Exam Preparation Guide May 2017 Fion Zhang/ Charlie Chong

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NACE MR0175-Carbon Steel Wi E Written Exam NACE-MR0175-Carbon Steel -001 Exam Preparation Guide May 2017 Exam Preparation Guide May 2017 Introduction The MR0175-Carbon Steel written exam is designed to assess whether a candidate has the g requisite knowledge and skills that a minimally qualified MR0175 Certified User- Carbon Steel must possess. The exam comprises 50 multiple-choice questions that are based on the MR0175 Standard Parts 1 and 2. multiple-choice p Fion Zhang/ Charlie Chong https://www.naceinstitute.org/uploadedFiles/Certification/Specialty_Program/MR0175-Carbon-Steel-EPG.pdf

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EXAM BOK EXAM BOK Suggested Study Material Suggested Study Material  NACE MR0175/ISO 15156 Standard 20171015-OK  EFC Publication 17  NACE TM0177  NACE TM0198 NACE TM0316 Books  Introductory Handbook for NACE MR0175 Fion Zhang/ Charlie Chong

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Fion Zhang/ Charlie Chong

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Fion Zhang/ Charlie Chong

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Reading 1: Reading 1: Fion Zhang/ Charlie Chong http://www.emerson.com/documents/automation/140706.pdf

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Reading 1: Reading 1: Sulfide Stress Cracking NACE MR0175 2002 MR0175/ISO 15156 --NACE MR0175-2002 MR0175/ISO 15156 Emerson Experiences Fion Zhang/ Charlie Chong http://www.emerson.com/documents/automation/140706.pdf

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The Details NACE MR0175 “Sulfide Stress Corrosion Cracking Resistant Metallic NACE MR0175 Sulfide Stress Corrosion Cracking Resistant Metallic Materials for Oil Field Equipment” is widely used throughout the world. In late 2003 it became NACE MR0175/ ISO 15156 “Petroleum and Natural Gas Industries - Materials for Use in H2S-Containing Environments in Oil and Gas Production.” These standards specify the proper materials heat treat conditions and strength levels required to provide good service life in sour gas and oil environments. NACE International formerly the National Association of Corrosion Engineers is a worldwide technical organization which studies various aspects of corrosion and the damage that may result in refineries various aspects of corrosion and the damage that may result in refineries chemical plants water systems and other types of industrial equipment. MR0175 was first issued in 1975 but the origin of the document dates to 1959 hfiiWtCdldthi ii when a group of engineers in Western Canada pooled their experience in successful handling of sour gas. The group organized as a NACE committee and in 1963 issued specification 1B163 “Recommendations of Materials or p Sour Service.” In 1965 NACE organized a nationwide committee which issued 1F166 in 1966 and MR0175 in 1975. Fion Zhang/ Charlie Chong

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Revisions were issued on an annual basis as new materials and processes were added Revisions had to receive unanimous approval from the were added. Revisions had to receive unanimous approval from the responsible NACE committee. Fion Zhang/ Charlie Chong

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In the mid-1990’s the European Federation of Corrosion EFC issued 2 reports closely related to MR0175 Publication 16 issued 2 reports closely related to MR0175 Publication 16 “Guidelines on Materials Requirements for Carbon and Low Alloy Steels for H2S-Containing Environments in Oil and Gas Production” Steels for H2S Containing Environments in Oil and Gas Production and Publication 17 “Corrosion Resistant Alloys for Oil and Gas Production: Guidance on General Requirements and Test Methods for H2S Service.” EFC is located in London England. Th I t ti l O i ti f St d di ti ISO i The International Organization for Standardization ISO is a worldwide federation of national standards bodies from more than 140 countries One organization from each country acts as the 140 countries. One organization from each country acts as the representative for all organizations in that country. The American National Standards Institute ANSI is the USA representative in ISO. Technical Committee 67 “Materials Equipment and Offshore Structures for Petroleum Petrochemical and Natural Gas I d t i ” t d th t NACE bl d th diff t i Industries” requested that NACE blend the different sour service documents into a single global standard. Fion Zhang/ Charlie Chong

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This task was completed in late 2003 and the document was issued as ISO standard NACE MR0175/ISO 15156 It is issued as ISO standard NACE MR0175/ISO 15156. It is now maintained by ISO/TC 67 Work Group 7 a 12-member “Maintenance Panel” and a 40-member Oversight Committee Maintenance Panel and a 40 member Oversight Committee under combined NACE/ISO control. The three committees are an international group of users manufacturers and service providers. Membership is approved by NACE and ISO based on technical knowledge and experience. Terms are limited. Previously some bthNACETkG hd df 25 members on the NACE Task Group had served for over 25 years. Fion Zhang/ Charlie Chong

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NACE MR0175/ISO 15156 is published in 3 volumes.  Part 1: General Principles for Selection of Cracking Resistant  Part 1: General Principles for Selection of Cracking-Resistant Materials  Part 2: Cracking-Resistant Carbon and Low Alloy Steels and the  Part 2: Cracking Resistant Carbon and Low Alloy Steels and the Use of Cast Irons  Part 3: Cracking-Resistant CRA’s Corrosion-Resistant Alloys and Other Alloys Fion Zhang/ Charlie Chong

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NACE MR0175/ISO 15156 applies only to petroleum production drilling gathering and flow line equipment and field processing drilling gathering and flow line equipment and field processing facilities to be used in H2S bearing hydrocarbon service. In the past MR0175 only addressed sulfide stress cracking SSC In past MR0175 only addressed sulfide stress cracking SSC. In NACE MR0175/ISO 15156 however but both SSC and chloride stress corrosion cracking SCC are considered. Question D SCC l d t th f Cl Does SCC always due to the presence of Cl - See http://oilfieldwiki.com/wiki/Stress_corrosion_cracking Fion Zhang/ Charlie Chong

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While clearly intended to be used only for oil field equipment industry has applied MR0175 in to many other areas including refineries LNG has applied MR0175 in to many other areas including refineries LNG plants pipelines and natural gas systems. The judicious use of the document in these applications is constructive and can help the document in these applications is constructive and can help prevent SSC failures wherever H2S is present. Saltwater wells and saltwater handling facilities are not covered by NACE MR0175/ ISO 15156. These are covered by NACE Standard RP0475 “Selection of Metallic Materials to Be Used in All Phases of Water Handling for I j ti i t OilBi F ti ” Injection into Oil-Bearing Formations.” Fion Zhang/ Charlie Chong

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When new restrictions are placed on materials in NACE MR0175/ ISO 15156 or when materials are deleted from this standard ISO 15156 or when materials are deleted from this standard materials in use at that time are in compliance. This includes materials listed in MR0175-2002 but not listed in NACE materials listed in MR0175 2002 but not listed in NACE MR0175/ISO 15156. However if this equipment is moved to a different location and exposed to different conditions the materials must be listed in the current revision. Alternatively successful use of materials outside the limitations of NACE MR0175/ISO 15156 b t t d b lifi ti t ti th t d d Th may be perpetuated by qualification testing per the standard. The user may replace materials in kind for existing wells or for new wells within a given field if the environmental conditions of the wells within a given field if the environmental conditions of the field have not changed. Fion Zhang/ Charlie Chong

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New Sulfide Stress Cracking Standard for Refineries DB hPi i lE i Mt il tE P Don Bush Principal Engineer - Materials at Emerson Process Management Fisher Valves is a member and former chair of a NACE task group that has written a document for refinery a NACE task group that has written a document for refinery applications NACE MR0103. The title is “Materials Resistant to Sulfide Stress Cracking in Corrosive Petroleum Refining gg Environments.” The requirements of this standard are very similar to the pre-2003 MR0175 for many materials. When applying this td dth h t ti k t il d ith standard there are changes to certain key materials compared with NACE MR0175-2002. Fion Zhang/ Charlie Chong

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Responsibility It has always been the responsibility of the end user to determine It has always been the responsibility of the end user to determine the operating conditions and to specify when NACE MR0175 applies This is now emphasized more strongly than ever in NACE applies. This is now emphasized more strongly than ever in NACE MR0175/ISO 15156.  The manufacturer is responsible for meeting the metallurgical requirements of NACE MR0175/ISO 15156.  It is the end user’s responsibility to ensure that a material will be ti f t i th i t d d i t satisfactory in the intended environment. Fion Zhang/ Charlie Chong

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Some of the operating conditions which must be considered include pressure temperature corrosiveness fluid properties etc When pressure temperature corrosiveness fluid properties etc. When bolting components are selected the pressure rating of flanges could be affected be affected. It is always the responsibility of the equipment user to convey the environmental conditions to the equipment supplier particularly if the equipment will be used in sour service. Th i ti f NACE MR0175/ISO 15156 th The various sections of NACE MR0175/ISO 15156 cover the commonly available forms of materials and alloy systems. The requirements for heat treatment hardness levels conditions of requirements for heat treatment hardness levels conditions of mechanical work and post-weld heat treatment are addressed for each form of material. Fabrication techniques bolting platings and coatings are also addressed. Fion Zhang/ Charlie Chong

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Applicability of NACE MR0175/ISO 15156 L t ti f H2S 0 05 i 0 3 kP H2S ti l Low concentrations of H2S 0.05 psi 03 kPa H2S partial pressure and low pressures 65 psia or 450 kPa are considered outside the scope of NACE MR0175/ISO 15156 The low stress outside the scope of NACE MR0175/ISO 15156. The low stress levels at low pressures or the inhibitive effects of oil may give satisfactory performance with standard commercial equipment. yp q p Many users however have elected to take a conservative approach and specify compliance to either NACE MR0175 or NACE MR0175/ISO 15156 ti bl t f H2S i MR0175/ISO 15156 any time a measurable amount of H2S is present. The decision to follow these specifications must be made by the user based on economic impact the safety aspects made by the user based on economic impact the safety aspects should a failure occur and past field experience. Legislation can impact the decision as well. Such jurisdictions include the Texas pj Railroad Commission and the U.S. Minerals Management Service offshore. The Alberta Canada Energy Conservation Board dfth ifiti recommends use of the specifications. Fion Zhang/ Charlie Chong

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Facilities operating at a total absolute pressure below 0 45 MPa 65 psi Facilities operating at a total absolute pressure below 0.45 MPa 65 psi Fion Zhang/ Charlie Chong

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Figure 1. Photomicrograph Showing Stress Corrosion Cracking Fion Zhang/ Charlie Chong

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Basics of Sulfide Stress Cracking SSC and Stress C i C ki SCC Corrosion Cracking SCC SSC and SCC are cracking processes that develop in the presence of water corrosion and surface tensile stress It is a progressive type of failure that corrosion and surface tensile stress. It is a progressive type of failure that produces cracking at stress levels that are well below the material’s tensile strength. The break or fracture appears brittle with no localized yielding lti df ti l ti Rth th i l k t k ffi plastic deformation or elongation. Rather than a single crack a network of fine feathery branched cracks will form see Figure 1. Pitting is frequently seen an will serve as a stress concentrator to initiate cracking. With SSC seen an will serve as a stress concentrator to initiate cracking. With SSC hydrogen ions are a product of the corrosion process Figure 2. These ions pick up electrons from the base material producing hydrogen atoms. At that it t h d t bi t f h d l l M t point two hydrogen atoms may combine to form a hydrogen molecule. Most molecules will eventually collect form hydrogen bubbles and float away harmlessly. However some percentage of the hydrogen atoms will diffuse into y p g yg the base metal and embrittle the crystalline structure. When a certain critical concentration of hydrogen is reached and combined with a tensile stress exceeding a threshold level SSC will occur H2S does not actively participate exceeding a threshold level SSC will occur. H2S does not actively participate in the SSC reaction however sulfides act to promote the entry of the hydrogen atoms into the base material. Fion Zhang/ Charlie Chong

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Figure 2. Schematic Showing the Generation of Hydrogen Producing SSC Fion Zhang/ Charlie Chong

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As little as 0.05 psi 03 kPa H2S partial pressure in 65 psia 450 kPa hydrocarbon gas can cause SSC of carbon and low alloy steels hydrocarbon gas can cause SSC of carbon and low alloy steels. Sulfide stress cracking is most severe at ambient temperature particularly in the range of 20 ° to 120 °F-6 ° to 49 °C Below 20 °F- particularly in the range of 20 to 120 F 6 to 49 C. Below 20 F 6 °C the diffusion rate of the hydrogen is so slo that the critical concentration is never reached. Above 120 °F 49 °C the diffusion rate is so fast that the hydrogen atoms pass through the material in such a rapid manner that the critical concentration is not reached. Key-Numbers: Key Numbers: ¤ 65 psia absolute pressure ¤ as litle as 0.05 psi 03 kPa H2S partial pressure Fion Zhang/ Charlie Chong

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Chloride SCC is widely encountered and has been extensively studied Much is still unknown however about its mechanism studied. Much is still unknown however about its mechanism. 1 One theory says that hydrogen generated by the corrosion 1. One theory says that hydrogen generated by the corrosion process diffuses into the base metal in the atomic form and embrittles the lattice structure. 2. A second more widely accepted theory proposes an lt h i l h i Stil t l d ith electrochemical mechanism. Stainless steels are covered with a protective chromium oxide film. The chloride ions rupture the film at weak spots resulting in anodic bare and cathodic film covered at weak spots resulting in anodic bare and cathodic film covered sites. The galvanic cell produces accelerated attack at the anodic sites which when combined with tensile stresses produces cracking. Aii i t ti i i dt d SCCA th A minimum ion concentration is required to produce SCC. As the concentration increases the environment becomes more severe reducing the time to failure reducing the time to failure. Fion Zhang/ Charlie Chong

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Temperature also is a factor in SCC. In general the likelihood of SCC increases with increasing temperature A minimum threshold temperature increases with increasing temperature. A minimum threshold temperature exists for most systems below which SCC is rare. Across industry the generally accepted minimum temperature for chloride SCC of the 300 SST’s is about 160 °F 71 °C. NACE MR0175/ISO 15156 has set a very conservative limit of 140 °F 60 °C due to the synergistic effects of the chlorides H2S and low pH values As the temperature increases above these values the time to low pH values. As the temperature increases above these values the time to failure will typically decrease. Resistance to chloride SCC increases with higher alloy materials. This is reflected in the environmental limits set by NACE MR0175/ISO 15156. Environmental limits progressively increase from 400 Series SST and ferritic Environmental limits progressively increase from 400 Series SST and ferritic SST to 300 Series highly alloy austenitic SST duplex SST nickel and cobalt base alloys. Fion Zhang/ Charlie Chong

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Carbon Steel Cb dl ll t l h t bl i t t SSC d SCC Carbon and low-alloy steels have acceptable resistance to SSC and SCC however their application is often limited by their low resistance to general corrosion. The processing of carbon and low alloy steels must be carefully pg y y controlled for good resistance to SSC and SCC. The hardness must be less than 22 HRC. If welding or significant cold working is done stress relief is required Although the base metal hardness of a carbon or alloy steel is less required. Although the base metal hardness of a carbon or alloy steel is less than 22 HRC areas of the heat affected zone HAZ will be harder. PWHT will eliminate these excessively hard areas. ASME SA216 Grades WCB and WCC and SAME SA105 are the most commonly used body materials It is Fisher’s policy to stress relieve all welded commonly used body materials. It is Fisher s policy to stress relieve all welded carbon steels that are supplied to NACE MR0175/ISO 15156. Fion Zhang/ Charlie Chong

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All carbon steel castings sold to NACE MR0175/ISO 15156 requirements are produced using one of the following processes: produced using one of the following processes: 1. In particular product lines where a large percentage of carbon steel assemblies are sold as NACE MR0175/ISO 15156 compliant castings are ordered from the foundry with a requirement that the castings be either normalized or stress relieved following all weld repairs major or minor normalized or stress relieved following all weld repairs major or minor. Any weld repairs performed either major or minor are subsequently stress relieved. 2. In product lines where only a small percentage of carbon steel products are ordered NACE MR0175/ISO 15156 compliant stock castings are stress relieved whether they are weld repaired by Emerson Process stress relieved whether they are weld repaired by Emerson Process Management or not. This eliminates the chance of a minor foundry weld repair going undetected and not being stress relieved. Fion Zhang/ Charlie Chong

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ASME SA352 grades LCB and LCC have the same composition as WCB and WCC respectively They are heat treated differently and impact tested at WCC respectively. They are heat treated differently and impact tested at - 50 °F -46 °C to ensure good toughness in low temperature service. LCB and LCC are used in locations where temperatures commonly drop below the - 20 °F -29 °C permitted for WCB and WCC. LCB and LCC castings are processed in the same manner as WCB and WCC when required to meet NACE MR0175/ ISO 15156 NACE MR0175/ ISO 15156. For carbon and low-alloy steels NACE MR0175/ISO 15156 imposes some changes in the requirements for the weld procedure qualification report PQR. All new PQR’s will meet these requirements however it will take several years for Emerson Process Management and our suppliers to complete this years for Emerson Process Management and our suppliers to complete this work. At this time we will require user approval to use HRC. Fion Zhang/ Charlie Chong

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Carbon and Low-Alloy Steel Welding Hardness Ri t Requirements  HV-10 HV-5 or Rockwell 15N.  HRC testing is acceptable if the design stresses are less than 67 of the  HRC testing is acceptable if the design stresses are less than 67 of the minimum specified yield strength and the PQR includes PWHT.  Other methods require user approval.  250 HV 70 6 HR15N i  250 HV or 70.6 HR15N maximum.  22 HRC maximum if approved by user. Fion Zhang/ Charlie Chong

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Low-Alloy Steel Welding Hardness Requirements  All of the above apply with the additional requirement of stress relieve at 1150 °F 621 °C minimum after welding. g All new PQR’s at Emerson Process Management and our foundries will require hardness testing with HV 10 HV 5 or Rockwell 15N and HRC The require hardness testing with HV-10 HV-5 or Rockwell 15N and HRC. The acceptable maximum hardness values will be 250 HV or 70.6 HR15N and 22 HRC. Hardness traverse locations are specified in NACE MR0175/ISO 15156 part 2 as a function of thickness and weld configuration. The number and locations of production hardness tests are still outside the scope of the standard The maximum allowable nickel content for carbon and low alloy standard. The maximum allowable nickel content for carbon and low-alloy steels and their weld deposits is 1. Note: A.2.1.6 Cold deformation and thermal stress relief Carbon and low-alloy steels shall be thermally stress-relieved following any cold deforming by rolling cold forging or other manufacturing process that results in a permanent outer fibre deformation greater than 5 . Thermal stress relief shall be manufacturing process that results in a permanent outer fibre deformation greater than 5 . Thermal stress relief shall be performed in accordance with an appropriate code or standard. The minimum stress-relief temperature shall be 595 °C 1100 °F. The final maximum hardness shall be 22 HRC except for pipe fittings made from ASTM A234 grade WPB or WPC for which the final hardness shall not exceed 197 HBW. Fion Zhang/ Charlie Chong

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Table A.1 — Maximum acceptable hardness values for carbon steel carbon manganese steel and carbon-manganese steel and low-alloy steel welds Fion Zhang/ Charlie Chong

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Low alloy steels like WC6 WC9 and C5 are acceptable to NACE MR0175/ISO 15156 to a maximum hardness of 22 HRC These castings must MR0175/ISO 15156 to a maximum hardness of 22 HRC. These castings must all be stress relieved to FMS 20B52 . The compositions of C12 C12a F9 and F91 materials do not fall within the definition of “low alloy steel” in NACE MR0175/ISO 15156 therefore these materials are not acceptable materials are not acceptable. A few customers have specified a maximum carbon equivalent CE for carbon steel. The primary driver for this requirement is to improve the SSC resistance in the as-welded condition. Fisher’s practice of stress relieving all carbon steel negates this need. Decreasing the CE reduces the hardenability of the steel and presumably i i t t lfid t ki SSC improves resistance to sulfide stress cracking SSC. Because reducing the CE decreases the strength of the steel there is a limit to how far the CE can be reduced. Fion Zhang/ Charlie Chong

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ANSI/NACE MR0175/ISO 15156-1 3.14 low-alloy steel steel with a total alloying element content of less than about 5 mass fraction but more than specified for carbon steel 3.3 3.3 carbon steel alloy of carbon and iron containing up to:  2 mass fraction carbon and  up to 1.65 mass fraction manganese and residual quantities of other  up to 1.65 mass fraction manganese and residual quantities of other elements except those intentionally added in specific quantities for deoxidation usually ili d/ liiNt1t t C b t l dith t l silicon and/or aluminium Note 1 to entry: Carbon steels used in the petroleum industry usually contain less than 0.8 mass fraction carbon. Fion Zhang/ Charlie Chong

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Cast Iron Gtiti dhit ti tb df Gray austenitic and white cast irons cannot be used for any pressure- retaining parts due to low ductility. Ferritic ductile iron to ASTM A395 is acceptable when permitted by ANSI API or other industry standards. pp y y Fion Zhang/ Charlie Chong

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Gray Cast Iron Fion Zhang/ Charlie Chong

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Gray Cast Iron Fion Zhang/ Charlie Chong

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White Cast Iron Fion Zhang/ Charlie Chong

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White Cast Iron Fion Zhang/ Charlie Chong

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Austenitic Cast Iron Fion Zhang/ Charlie Chong

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Austenitic Cast Iron Fion Zhang/ Charlie Chong

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Austenitic Cast Iron Fion Zhang/ Charlie Chong

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Meallable Cast Iron Fion Zhang/ Charlie Chong

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Meallable Cast Iron Fion Zhang/ Charlie Chong

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Meallable Cast Iron Fion Zhang/ Charlie Chong

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Fe-C Equillibrium Equillibrium Diagram Fion Zhang/ Charlie Chong

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Optional Reading Non CS Scope - Non CS Scope Fion Zhang/ Charlie Chong

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Stainless Steel 400 S i St i l St l 400 Series Stainless Steel UNS 410 410 SST CA15 cast 410 420 420 SST and several other martensitic grades must be double tempered to a maximum other martensitic grades must be double tempered to a maximum hardness of 22 HRC. PWHT is also required. An environmental limit now applies to the martensitic grades 1.5 psi 10 kPa H2S ti l d H t th l t 3 5 416 416 SST partial pressure and pH greater than or equal to 3.5 416 416 SST is similar to 410 410 with the exception of a sulfur addition to produce free machining characteristics. Use of 416 and other free produce free machining characteristics. Use of 416 and other free machining steels is not permitted by NACE MR0175/ISO 15156. CA6NM is a modified version of the cast 410 stainless steel. NACE MR0175/ISO 15156 ll it b t ifi th NACE MR0175/ISO 15156 allows its use but specifies the exact heat treatment required. Generally the carbon content must be restricted to 0.03 maximum to meet the 23 HRC maximum hardness. PWHT is required for CA6NM. The same environmental limit applies 1.5 psi 10 kPa H2S partial pressure and pH greater than or equal to 3 5 and pH greater than or equal to 3.5. Fion Zhang/ Charlie Chong

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300 Series Stainless Steel S l h h b d ith th i t f th Several changes have been made with the requirements of the austenitic 300 Series stainless steels. Individual alloys are no longer listed. All alloys with the following elemental ranges gy g g are acceptable: C 0.08 maximum Cr 16 minimum Ni 8 minimum P 0.045 maximum S 0.04 maximum Mn 2.0 maximum and Si 2 0 maximum Other alloying elements are maximum and Si 2.0 maximum. Other alloying elements are permitted. The other requirements remain solution heat treated condition 22 HRC maximum and free of cold work designed to improve mechanical properties. The cast and wrought equivalents of 302 304 CF8 S30403 CF3 310 CK20 316 CF8M S31603 CF3M 317 CG8M S31703 CG3M 321 347 CF8C and CF3M 317 CG8M S31703 CG3M 321 347 CF8C and N08020 CN7M are all acceptable per NACE MR0175/ISO 15156. Fion Zhang/ Charlie Chong

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Environmental restrictions now apply to the 300 Series SST. The limits are 15 psia 100 kPa H2S partial pressure a maximum limits are 15 psia 100 kPa H2S partial pressure a maximum temperature of 140 °F 60 °C and no elemental sulfur. If the chloride content is less than 50 mg/L 50 ppm the H2S partial pressure must be less than 50 psia 350 kPa but there is no temperature limit. There is less of a restriction on 300 Series SST in oil and gas processing and injection facilities. If the chloride content in aqueous solutions is low typically less than 50 mg/L or 50 ppm chloride in operations after separation there are no limits for austenitic stainless steels highly alloyed austenitic stainless steels austenitic stainless steels highly alloyed austenitic stainless steels duplex stainless steels or nickel-based alloys. di --------------- more reading Fion Zhang/ Charlie Chong

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Reading 2 Reading 2 Stress Corrosion Cracking http://oilfieldwiki com/wiki/Stress corrosion cracking http://oilfieldwiki.com/wiki/Stress_corrosion_cracking Fion Zhang/ Charlie Chong

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SCC St i ki SCC i th th f k i i Stress corrosion cracking SCC is the growth of cracks in a corrosive environment. It can lead to unexpected sudden failure of normally ductile metals subjected to a tensile stress especially at elevated temperature in the jpy p case of metals. SCC is highly chemically specific in that certain alloys are likely to undergo SCC only when exposed to a small number of chemical environments The chemical environment that causes SCC for a given alloy is environments. The chemical environment that causes SCC for a given alloy is often one which is only mildly corrosive to the metal otherwise. Hence metal parts with severe SCC can appear bright and shiny while being filled with microscopic cracks. This factor makes it common for SCC to go undetected prior to failure. SCC often progresses rapidly and is more common among alloys than pure metals The specific environment is of crucial importance and alloys than pure metals. The specific environment is of crucial importance and only very small concentrations of certain highly active chemicals are needed to produce catastrophic cracking often leading to devastating and unexpected failure.1 The stresses can be the result of the crevice loads due to stress The stresses can be the result of the crevice loads due to stress concentration or can be caused by the type of assembly or residual stresses from fabrication e.g. cold working the residual stresses can be relieved by li annealing. Fion Zhang/ Charlie Chong

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Metals attacked  C t i t iti til t l dl i i ll k i th  Certain austenitic stainless steels and aluminium alloys crack in the presence of chlorides  mild steel cracks in the presence of alkali boiler cracking and nitrates p g  copper alloys crack in ammoniacal solutions season cracking. This limits the usefulness of austenitic stainless steel for containing water with This limits the usefulness of austenitic stainless steel for containing water with higher than few ppm content of chlorides at temperatures above 50 °C. Worse still high-tensile structural steels crack in an unexpectedly brittle manner in a whole variety of aqueous environments especially containing chlorides. With the possible exception of the latter which is a special example of hydrogen cracking all the others display the phenomenon of subcritical crack growth cracking all the others display the phenomenon of subcritical crack growth i.e. small surface flaws propagate usually smoothly under conditions where fracture mechanics predicts that failure should not occur. That is in the presence of a corrodent cracks develop and propagate well below K Ic . In fact the subcritical value of the stress intensity designated as K Iscc may be less than 1 of K I as the following table shows: than 1 of K Ic as the following table shows: Fion Zhang/ Charlie Chong

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K IC K ISCC Fion Zhang/ Charlie Chong

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K IC K ISCC Fion Zhang/ Charlie Chong

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Crack growth Th b iti l t f ti b tt ib t d t th h i l The subcritical nature of propagation may be attributed to the chemical energy released as the crack propagates. That is elastic energy released + chemical energy surface energy + deformation energy. gy gy gy The crack initiates at K Iscc and thereafter propagates at a rate governed by the slowest process which most of the time is the rate at which corrosive ions can slowest process which most of the time is the rate at which corrosive ions can diffuse to the crack tip. As the crack advances so K rises because crack length appears in the calculation of stress intensity. Finally it reaches K Ic whereupon fast fracture ensues and the component fails ensues and the component fails. One of the practical difficulties with SCC is its unexpected nature. Stainless steels for example are employed because under most conditions they are "passive" i.e. effectively inert. Very often one finds a single crack has propagated while the rest of the metal surface stays apparently unaffected propagated while the rest of the metal surface stays apparently unaffected. The crack propagates perpendicular to the applied stress. Fion Zhang/ Charlie Chong

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Prevention SCC i th lt f bi ti f th f t SCC is the result of a combination of three factors –  a susceptible material  exposure to a corrosive environment and p  tensile stresses above a threshold. If you eliminate any one of these factors SCC initiation becomes impossible If you eliminate any one of these factors SCC initiation becomes impossible. The conventional approach to controlling the problem has been to develop new alloys that are more resistant to SCC. This is a costly proposition and can require a massive time investment to achieve only marginal success. Fion Zhang/ Charlie Chong

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Examples A 32 i h di t t i i i li th f N t hit h A 32 inch diameter gas transmission pipeline north of Natchitoches Louisiana belonging to the Tennessee Gas Pipeline exploded and burned from SCC on March 4 1965 killing 17 people. At least 9 others were injured g pp j and 7 homes 450 feet from the rupture were destroyed.23 SCC caused the catastrophic collapse of the Silver Bridge in December 1967 SCC caused the catastrophic collapse of the Silver Bridge in December 1967 when an eyebar suspension bridge across the Ohio river at Point Pleasant West Virginia suddenly failed. The main chain joint failed and the whole structure fell into the river killing 46 people in vehicles on the bridge at the time. Rust in the eyebar joint had caused a stress corrosion crack which went critical as a result of high bridge loading and low temperature The failure was critical as a result of high bridge loading and low temperature. The failure was exacerbated to increase the severity bitterness or violence of disease ill feeling etc. aggravate by a high level of residual stress in the eyebar. The disaster led to a nationwide reappraisal of bridges.4 Fion Zhang/ Charlie Chong

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A classic example of SCC is season cracking of brass cartridge cases a problem experienced by the British army in India in the early 19th century It problem experienced by the British army in India in the early 19th century. It was initiated by ammonia from dung and horse manure decomposing at the higher temperatures of the spring and summer. There was substantial residual stress in the cartridge shells as a result of cold forming. The problem was solved by annealing the shells to ameliorate the stress. Fion Zhang/ Charlie Chong

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The collapsed Silver Bridge as seen from the Ohio side Fion Zhang/ Charlie Chong

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The collapsed Silver Bridge as seen from the Ohio side Fion Zhang/ Charlie Chong

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The collapsed Silver Bridge as seen from the Ohio side Fion Zhang/ Charlie Chong

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SCC Aluminum- Ammonia NH3 Fion Zhang/ Charlie Chong

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SCC Aluminum- Ammonia NH3 Fion Zhang/ Charlie Chong

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Reading 3 Reading 3 Corrosion problems in production The WikiPetrol The WikiPetrol Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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Corrosion problems in production Ci f tlith f ti bl Corrosion of metal in the presence of water is a common problem across many industries. The fact that most oil and gas production includes co- produced water makes corrosion a pervasive issue across the industry. Age pp yg and presence of corrosive materials such as carbon dioxide CO2 and hydrogen sulfide H2S exacerbate the problem. Corrosion control in oil and gas production is reviewed in depth in Treseder and Tuttle1 Brondel et al.2 and NACE3 from which some of the following material is abstracted. Fion Zhang/ Charlie Chong

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Corrosion chemistry of steels I i i h tl th d i ll ffii tl ti t t Iron is inherently thermodynamically sufficiently active to react spontaneously with water corrosion generating soluble iron ions and hydrogen gas. The utility of iron alloys depends on minimizing the corrosion yg g y y p g rate. Corrosion of steel is an “electrochemical process” involving the transfer of electrons from iron atoms in the metal to hydrogen ions or oxygen in water. The corrosion reaction of iron with H2S under anearobic condition is described The corrosion reaction of iron with H2S under anearobic condition is described by the equation Corrosion under anaerobic condition commonly found in oilfield production: Anode: Fe + H O → Fe 2+ +2e - +H O Anode: Fe + H 2 O → Fe 2 + 2e + H 2 O Cathode: H 2 S + H 2 O → H + + HS - + H 2 O HS - + H 2 O → H + + S 2- + H 2 O Fion Zhang/ Charlie Chong

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This separation of the overall corrosion process into two reactions is not an electrochemical nuance these processes generally do take place at separate electrochemical nuance these processes generally do take place at separate locations on the same piece of metal. This separation requires the presence of a medium to complete the electrical circuit between anode site of iron dissolution and cathode site for corrodant reduction. Electrons travel in the metal phase but the ions involved in the corrosion process cannot. Ions require the presence of water hence corrosion requires the presence of require the presence of water hence corrosion requires the presence of water. This overall process is shown schematically in Fig. 1. 3 The space between the anode and cathode may be small or large depending on a number of factors. Fion Zhang/ Charlie Chong

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Acid is not the only corrodant possible. Another common cathodic process is the reduction of oxygen which is written as : the reduction of oxygen which is written as : O 2 + 4H + + 4e - → 2H 2 O This reaction can also take place at a location different from that of iron This reaction can also take place at a location different from that of iron dissolution. The other chemical constituents in the vicinity of the anodic sites determine the ultimate chemical fate of the Fe ++ ion such as the precipitation of iron- containing solids on or near the corroding surface. containing solids on or near the corroding surface. The net rate of corrosion is determined by how fast the corrodant arrives at the iron-atom/water interface how much corrodant is present the electrical t til fth d t h hih t tilth d potential energy of the corrodant oxygen has a higher potential than do protons and the intrinsic rate of the cathodic reactions—electron transfer processes involving protons and oxygen are not instantaneous and depend on pgp yg p the nature of the solid surface on which they occur. Fion Zhang/ Charlie Chong

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“How fast the corrodants arrive” has two aspects:  Mass transport in the corroding fluid  Mass transport in the corroding fluid  Permeating surface barriers between the iron metal and the water phase Surface barriers are placed barriers such as:  Paint or plastic coatings  Passivating oxide films inherent to the metal discussed later  Passivating oxide films inherent to the metal discussed later  Low-permeability corrosion products e.g. siderite as formed in the presence of certain oils and/or inhibitors http://petrowiki.org/Corrosion_problems_in _production Fion Zhang/ Charlie Chong

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Nature of steels All i i ith b ll 0 2 t 1 f t l l ll t l Alloying iron with carbon usually 0.2 to 1 forms steel low-alloy steel—a far stronger metal than iron hence suitable for oilfield use. Other components can be added to iron to enhance corrosion-resistance properties. pp Some of the carbon added is insoluble forming iron carbide Fe3C which accelerates the cathodic processes necessary for corrosion to take place accelerating the corrosion rate One of the major ubiquitous impurities in steel accelerating the corrosion rate. One of the major ubiquitous impurities in steel is sulfur and it is a major source of corrosion instability. This element is highly insoluble in iron and precipitates in the form of insoluble sulfide inclusions in particular MnS and Mn FeS. These inclusions are generally the sites of pitting discussed later. 4 Grain boundaries are also areas that are chemically active 3 When iron Grain boundaries are also areas that are chemically active. 3 When iron solidifies during casting the atoms which are randomly distributed in the liquid state arrange themselves in a crystalline array. This ordering usually begins simultaneously at many points in the liquid and as these blocks of crystals and grains meet there is a mismatch in the boundaries. There are areas of higher energy Chemical impurities in the melt tend to accumulate at areas of higher energy. Chemical impurities in the melt tend to accumulate at these grain boundaries and are more susceptible to chemical attack than the iron surface itself. Fion Zhang/ Charlie Chong

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Plain carbon steels are processed by one of four heat treatments:  Annealing  Annealing  Normalizing  Spherodizing  Quench and tempering These treatments determine in part the physical and corrosion properties of These treatments determine in part the physical and corrosion properties of the metal. Annealing or normalizing results in greater corrosion resistance than spherodizing or quench and tempering. The logic is that these treatments determine in large part of the physical dimensions and distribution of the impurities and inclusions in the metal. Fion Zhang/ Charlie Chong

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The corrosion products formed in oxygen-containing water on mild steel are FeOOH likely amorphous and magnetite 4 Below 200 °C these oxides in the FeOOH likely amorphous and magnetite. 4 Below 200 C these oxides in the absence of reactive inclusions are protective. In the presence of dissolved CO 2 FeCO 3 films form which can sometimes be protecting discussed later. The compositions of corrosion-resistant alloys CRAs are chosen to spontaneously generate surface oxide films that will be stable and spontaneously generate surface oxide films that will be stable and impermeable in the presence of the more aggressive corrodants. In oilfield use it is also required that these films spontaneously reform if ruptured as for example during and after erosion by sand or scratching by wireline/caliper tools. CRAs include the ferrous stainless steels and nonferrous nickel and cobalt alloys. Stainless steels contain at least 12 chromium. These alloys cobalt alloys. Stainless steels contain at least 12 chromium. These alloys passivate in oxidizing environments through the formation of a thin layer of chromium oxide—containing film on the surface of the alloy. The crystallinity of thi fil d ith i i C t t i th t l b i this film decreases with increasing Cr content in the steel becoming more glass-like and more protective. 5 Again various inclusions can be weak points in the passivating film. The surfaces of nickel-based CRAs such as Incoloy pg y 800™ are a passivating nickel ferrite Ni0.8Fe2.2O4. Fion Zhang/ Charlie Chong

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There are four classes of stainless steels that are based on chemical content metallurgical structure and mechanical properties These classes are: metallurgical structure and mechanical properties. These classes are:  Martensitic  Ferritic  Austenitic  Duplex  PH  PH The manufacturing processes for CRAs are more complex than those producing low-alloy steels. Stainless steels are less costly than the nickel and cobalt alloys though they are 1.5 to 20 times more expensive than low-alloy steels. Fion Zhang/ Charlie Chong

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Oilfield corrosion Oilfi ld i b di id d i t i d t " t" Oilfield corrosion can be divided into corrosions due to oxygen "sweet" corrosion and "sour" corrosion. Corrosion because of oxygen is found with surface equipment and can be found downhole with the oxygen introduced by waterflooding pressure maintenance gas lifting or completion and/or workover fluids It is the major maintenance gas lifting or completion and/or workover fluids. It is the major corrodant of offshore platforms at and below the tide line. The chemistry of this process follows the equations given below Oxidation of iron at points of stress in the crystal lattice: 2Fes 2Fe 2+ aq + 4e - 2Fes 2Fe 2 aq + 4e Reduction of water at the site of carbon impurities: O 2 g + H 2 Ol + 4e - 4OH - aq Overall equation: Overall equation: 2Fes + O 2 g + H 2 Ol FeOH 2 . Fion Zhang/ Charlie Chong http://www.chemicalformula.org/chemistry-help/corrosion

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The ironII hydroxide is converted to rust through a serious of reactions. The ironII hydroxide firstly oxides to ironIII oxide. 1. FeOH 2 s → FeOH 3 oxidation The ironIII oxide then changes to rust through a dehydration reaction. 2. FeOH 3 s → Fe 2 O 3 .nH 2 Os or rust dehydration 2FeOH 3 → Fe 2 O 3 + 3H 2 O Rt dh l l t th f fth tlThi th tlt Rust adheres loosely to the surface of the metal. This exposes the metal to more and more water and oxygen allowing the process of rusting to continue. Fion Zhang/ Charlie Chong http://www.chemicalformula.org/chemistry-help/corrosion

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“Sweet” corrosion is generally characterized first by simple metal dissolution followed by pitting The corrodant is H+ derived from carbonic acid H2CO3 followed by pitting. The corrodant is H+ derived from carbonic acid H2CO3 and the dissolution of CO2 in the produced brine. The pitting leaves distinctive patterns e.g. “mesa” corrosion attributable to the metallurgical processing used in manufacturing the tubing. “Ringworm” corrosion is caused when welding is not followed by full-length normalizing of the tubular after processing Corrosion inhibitors and CRAs are effective in mitigating sweet processing. Corrosion inhibitors and CRAs are effective in mitigating sweet corrosion. Naphthenic acids and simple organic acids indigenous to crude oil also contribute to corrosion Fion Zhang/ Charlie Chong

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Sweet Corrosion The deterioration of metal due to contact with carbon dioxide or similar The deterioration of metal due to contact with carbon dioxide or similar corrosive agents but excluding hydrogen sulfide H 2 S. Sweet corrosion typically results in pitting or material loss and occurs where steel is exposed to carbon dioxide and moisture http://www.glossary.oilfield.slb.com/en/Terms/s/sweet_corrosion.aspx Fion Zhang/ Charlie Chong

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Sweet Sweet The terms sweet and sour are a ftth lf tt reference to the sulfur content of crude oil. Early prospectors yp p would taste oil to determine its quality with low sulfur oil quality with low sulfur oil actually tasting sweet. Crude is currently considered sweet if it contains less than 0 5 sulfur contains less than 0.5 sulfur. Fion Zhang/ Charlie Chong

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“Sour” corrosion H 2 S results in the formation of various insoluble iron sulfides on the metal surface Not only is H S an acidic corrodant it also acts as a on the metal surface. Not only is H 2 S an acidic corrodant it also acts as a catalyst for both the anodic and cathodic halves of the corrosion reaction. Galvanic corrosion bimetallic corrosion is caused by the coupling of a corrosive and noncorrosive metal in the presence of a corrodant. Erosion is yet another category of corrosion Erosion corrosion is the acceleration of yet another category of corrosion. Erosion corrosion is the acceleration of corrosion because of the abrasion of metal surfaces by particulates e.g. sand. Finally there is corrosion caused by acids—those used to stimulate wells HCl and HF. Fion Zhang/ Charlie Chong

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Sweet The terms sweet and sour are a reference to the sulfur content of crude oil The terms sweet and sour are a reference to the sulfur content of crude oil. Early prospectors would taste oil to determine its quality with low sulfur oil actually tasting sweet. Crude is currently considered sweet if it contains less than 0.5 sulfur. Sweet crude is easier to refine and safer to extract and transport than sour Sweet crude is easier to refine and safer to extract and transport than sour crude. Because sulfur is corrosive light crude also causes less damage to refineries and thus results in lower maintenance costs over time. Due to all these factors sweet crude commands up to a 15 dollar premium per barrel over sour. Major locations where sweet crude is found include the Appalachian Basin in Eastern North America Western Texas the Bakken Formation of North Dkt dS kth th NthS fE NthAfi A t li Dakota and Saskatchewan the North Sea of Europe North Africa Australia and the Far East including Indonesia. Fion Zhang/ Charlie Chong http://www.petroleum.co.uk/sweet-vs-sour

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Sour Sour crude oil will have greater than 0 5 sulfur and some of this will be in the Sour crude oil will have greater than 0.5 sulfur and some of this will be in the form of hydrogen sulfide. Sour crude also contains more carbon dioxide. Most sulfur in crude is actually bonded to carbon atoms nevertheless high quantities of hydrogen sulfide in sour crude can pose serious health problems or even be fatal. Hydrogen sulfide is famous for its “rotten egg” smell which is only noticed at low concentrations. At moderate concentrations hydrogen sulfide can cause respiratory and nerve damage. At high concentrations it is instantly fatal. Exposure to high levels of hydrogen sulfide is thought to be in part responsible for Gulf War Syndrome which is characterized by chronic fatigue headaches for Gulf War Syndrome which is characterized by chronic fatigue headaches dizziness memory problems serious breathing problems and even birth defects. Hydrogen sulfide is so much of a risk that sour crude has to be t bili d i l f h d lfid b f it b t t d b il stabilized via removal of hydrogen sulfide before it can be transported by oil tankers. Sour crude is more common in the Gulf of Mexico Mexico South America and Canada. Crude produced by OPEC Member Nations also tends to be relatively sour with an average sulfur content of 1.77. Fion Zhang/ Charlie Chong http://www.petroleum.co.uk/sweet-vs-sour

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Gulf War Exposure to high levels of hydrogen sulfide is thought to be in part responsible Exposure to high levels of hydrogen sulfide is thought to be in part responsible for Gulf War Syndrome Fion Zhang/ Charlie Chong

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Gulf War Exposure to high levels of hydrogen sulfide is thought to be in part responsible Exposure to high levels of hydrogen sulfide is thought to be in part responsible for Gulf War Syndrome Fion Zhang/ Charlie Chong

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Gulf War Exposure to high levels of hydrogen sulfide is thought to be in part responsible Exposure to high levels of hydrogen sulfide is thought to be in part responsible for Gulf War Syndrome Fion Zhang/ Charlie Chong

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Gulf War Exposure to high levels of hydrogen sulfide is thought to be in part responsible Exposure to high levels of hydrogen sulfide is thought to be in part responsible for Gulf War Syndrome Fion Zhang/ Charlie Chong

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Gulf War Exposure to high levels of hydrogen sulfide is thought to be in part responsible Exposure to high levels of hydrogen sulfide is thought to be in part responsible for Gulf War Syndrome Fion Zhang/ Charlie Chong

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Gulf War Exposure to high levels of hydrogen sulfide is thought to be in part responsible Exposure to high levels of hydrogen sulfide is thought to be in part responsible for Gulf War Syndrome Fion Zhang/ Charlie Chong

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Oilfield corrosion can take specific forms: 1. Metal wastage 2. Pitting 3. Crevice corrosion 4. Intergranular corrosion 5 Stress corrosion cracking SCC 5. Stress corrosion cracking SCC 6. Blistering HIC 7. Embrittlement SCC/SSC... 8. Sulfide stress cracking SSC 9. Corrosion fatigue The first five forms involve primarily carbonic acid and/or dissolved oxygen as corrodants. Items 6 through 8 are induced primarily by H 2 S. Fion Zhang/ Charlie Chong

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Corrosive failure by uniform loss of metal is only infrequently seen during the production of oil and gas It is however the first step in corrosive failure of production of oil and gas. It is however the first step in corrosive failure of steels by means of localized corrosion. A circumstance for severe metal wastage is the pumping of poorly inhibited matrix stimulation acids. Pitting is the common failure mode of sweet corrosion and corrosion because of dissolved oxygen All passivating/protecting films on steel contain weak of dissolved oxygen. All passivating/protecting films on steel contain weak spots that will preferentially dissolve and form pits. As mentioned these areas are generally the sulfide inclusions. Chloride ion weakens the repassivating film allowing continued dissolution. The decreasing pH within the pit also enhances continued corrosion. The driver for theses processes is the large cathodic area of the metal oxide surface vs. the small anodic pit. Pitting is cathodic area of the metal oxide surface vs. the small anodic pit. Pitting is particularly dangerous because penetration through a tubular can occur relatively fast. Other corrosion mechanisms such as SCC frequently start at it O till dt thi i tt tt pits. Oxygen scavengers are typically used to remove this gas in an attempt to minimize the pitting problem. However small amounts may remain e.g. 20 ppb and these can be sufficient to induce corrosion. pp Fion Zhang/ Charlie Chong

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Carbonic acid the driver for sweet corrosion is a weak acid. The pH of the formation water depends on the CO partial pressure temperature and formation water depends on the CO 2 partial pressure temperature and alkalinity controlled primarily but not exclusively by the presence or absence of carbonate minerals in the formation. Shown in Fig. 2 as a function of CO 2 partial pressure are computed pH values for a seawater brine containing 140 ppm alkalinity and a seawater brine saturated in calcite at 50 and 150 °C substantially higher alkalinities For the common case of carbonate- substantially higher alkalinities. For the common case of carbonate containing reservoirs and moderate temperatures produced waters should have pH values of 6 or greater. Waters exposed to greater amounts of CO 2 in noncarbonate-containing reservoirs can have pH values of 4 or less. Fion Zhang/ Charlie Chong

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Fig. 2—Computed pH vs. pressure for a seawater brine exposed to a gas phase containing CO data are shown for seawater alone at 50 o Candfor phase containing CO 2 data are shown for seawater alone at 50 o C and for calcite-saturated seawater at 50 and 150 o C. Fion Zhang/ Charlie Chong

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Such corrosion induced by CO 2 is a function not only of CO 2 partial pressure and temperature but also of the crude oil Crude oil contains surface active and temperature but also of the crude oil. Crude oil contains surface-active chemicals—some oils contain more than others. These chemicals e.g. resins and asphaltenes can impact the corrosion process at least for low-alloy steels. For a fixed brine composition WORwater to oil ratio temperature and pressure corrosion in the presence of some crudes can be negligible while in the presence of others it can be extreme under identical while in the presence of others it can be extreme under identical environmental conditions. 67 Sweet corrosion generally results in the deposition of insoluble FeCO 3 siderite on the steel surface. It has been suggested that this selectivity to oil composition relates to the physical morphology of the FeCO 3 corrosion product—a compact tight film can protect the steel a loose poorly adherent film does not. 7 An example is shown in the steel a loose poorly adherent film does not. An example is shown in Fig. 3. The average uniform corrosion rate for steel in Crude B was 0.6 mil/yr the corrosion rate in Crude E was 26 mil/yr. Many corrosion inhibitors tl t b th h i i th ti f id it fil apparently act by the same mechanism i.e. the generation of siderite films similar and/or more compact than those formed from Crude B. Fion Zhang/ Charlie Chong

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Fig. 3—Scanning-electron-microscope micrographs X10K of the surface of the N 80 steel coupons after a 24 hour exposure at 186 o F to brine and 760 psi the N-80 steel coupons after a 24-hour exposure at 186 o F to brine and 760-psi CO 2 without crude oil upper left with 95 vol crude oil E upper right with 95 vol crude oil F lower left and with 95 vol crude oil B lower right all deposits are siderite courtesy of the Electrochemical Society. Fion Zhang/ Charlie Chong

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Fion Zhang/ Charlie Chong

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Alternatively it has been suggested that wettability plays the dominant role whereby the surface active components in the crude oil provide for a water whereby the surface-active components in the crude oil provide for a water- wet surface high corrosion rates or an oil-wet surface low corrosion rates. 8 Regardless of the mechanism crude oil can modify the corrosion rate. The penalty for ignoring the effect of crude-oil chemistry is the cost of overtreating or using more expensive alloys than are required. A crevice such as the junction space under a bolt or the physical junction of A crevice such as the junction space under a bolt or the physical junction of two metal parts is in effect a pit. Uniform corrosion can initiate in the presence of a corrodant within the crevice and continue driven by the large cathodic area outside the pit or crevice. Fion Zhang/ Charlie Chong

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Stress corrosion cracking is intergranular corrosion but it takes place only when the metal is under stress and in the presence of a corrodant The when the metal is under stress and in the presence of a corrodant. The corrodant can be specific—not all corrodants induce SCC on all alloys. Metal wastage is generally small SCC is often preceded by pitting. High-strength steels are more susceptible to SCC than low-strength alloys. The severity of intergranular corrosion generally depends on the metallurgical history of the steel Austenitic steels common stainless steels are particularly susceptible steel. Austenitic steels common stainless steels are particularly susceptible to intergranular attack. Blistering as well as embrittlement and sulfide stress cracking a subclass of SCC all stem from the same cause: the presence of H 2 S in the system and at the metal surface . The roots of the problem are in the mechanism for the the metal surface . The roots of the problem are in the mechanism for the cathodic discharge of hydrogen. The mechanism already discussed for the cathodic portion of the acid-induced corrosion process itself involves two t steps. H + + e - → H 1 H + H → H 2 2 Fion Zhang/ Charlie Chong

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i.e. the proton is first reduced to a hydrogen atom on the metal surface H • followed by the combination of two hydrogen atoms to yield hydrogen gas followed by the combination of two hydrogen atoms to yield hydrogen gas. Hydrogen sulfide inhibits the combination of hydrogen atoms as does arsenic HCN and some other corrosion inhibitors. Accordingly the hydrogen atoms can penetrate into the metal where they cause the corrosion problems that were already listed. This is shown schematically in Fig. 4 Fion Zhang/ Charlie Chong

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Fig. 4 —The alternatives for hydrogen atoms formed by the corrosion process: combination in the water phase to make gas diffusion into the metal to make combination in the water phase to make gas diffusion into the metal to make gas or embrittle steel penetration through the metal recombining to make gas a phenomenon also used to measure corrosionafter Schlumberger Oilfield Review. Fion Zhang/ Charlie Chong

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This hydrogen entry into low-strength steels can result in hydrogen blisters if there is a macroscopic defect in the steel such as an inclusion Such a void there is a macroscopic defect in the steel such as an inclusion. Such a void can provide a space for the hydrogen atoms to form hydrogen gas. Pressure builds and blisters form resulting in rupture and leakage. Embrittlement hydrogen-induced cracking and hydrogen embrittlement cracking causes failure at stresses well below the yield strength This cracking causes failure at stresses well below the yield strength . This phenomenon usually occurs only with high-strength hard steels generally those having yield strengths of 90000 psi or higher. Tubing and line pipe electric welded and seamless are susceptible to this effect. The dominating factor is the metallurgical structure of the steel relating to its method of manufacture. manufacture. Fion Zhang/ Charlie Chong

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Hydrogen Blister- Low strength steel Fion Zhang/ Charlie Chong

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Hydrogen Blister- Low strength steel Fion Zhang/ Charlie Chong

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Hydrogen Blister- Low strength steel Fion Zhang/ Charlie Chong

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Hydrogen Blister- Low strength steel Mi h h i h d b ittl t Micrography showing hydrogen embrittlement Fion Zhang/ Charlie Chong https://eduardodseng.wordpress.com/2013/11/10/different-types-of-corrosion/

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SSC cracking failure requires only low concentrations of H 2 S. The time to failure decreases as stress increases Cracking tendency increases as pH failure decreases as stress increases. Cracking tendency increases as pH decreases. SSC can be thought of in the same language as that used in describing hydraulic fracturing. There is a critical “stress intensity factor” below that at which a fracture crack will not propagate. This factor is related linearly to tensile strength. Some of this problem has been attributed to the effects of cold working on the alloys Alloys that were stress relieved were found to cold working on the alloys. Alloys that were stress relieved were found to increase in resistance to SSC. 9 Wells producing hydrocarbon liquids with the hydrogen sulfide are less susceptible to SSC pitting and weight loss. For example certain Canadian condensate wells have produced fluids with 40 mol H 2 Sand10CO 2 for 30 condensate wells have produced fluids with 40 mol H 2 S and 10 CO 2 for 30 years without serious corrosion problems. Stability is associated with a protective iron sulfide film wetted by the oil/liquid hydrocarbon. These wells lhd BHT f90 °C i lfid fil l ff ti i ti also had a BHT of 90 °C iron sulfide films are less effective in preventing corrosion above 110 °C. Fion Zhang/ Charlie Chong

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Canadian Condensate Wells Fion Zhang/ Charlie Chong

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Steels repeatedly stressed in a cyclical manner may fail in time corrosion fatigue It is required for failure that the stress be above a critical value called fatigue. It is required for failure that the stress be above a critical value called the “endurance limit” nominally 40 to 60 less than the tensile strength. The presence of a corrodant substantially reduces the fatigue life of a metal. Cyclic stress can be looked upon as a method of accelerating failure because of the other mechanisms previously described. Bimetallic corrosion/galvanic corrosion can occur when two metals are coupled in electrical contact and a corrodant is present. The more reactive metal corrodes faster while the less-reactive metal shows little or no corrosion. The more-reactive metal cathodically protects the less-reactive metal exploiting cathodic protection to prevent corrosion is discussed later. In exploiting cathodic protection to prevent corrosion is discussed later. In general the total corrosion of the anodic material is proportional to the exposed area of the cathodic material. Thus steel rivets in monel corrode very idl hil l i t i t l littl d rapidly while monel rivets in steel cause little damage. Fion Zhang/ Charlie Chong

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Anode to Cathode Ratio Fion Zhang/ Charlie Chong

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Weld-related corrosion is a variant of galvanic corrosion. When a metal is welded the welding process can generate a microstructure different from that welded the welding process can generate a microstructure different from that of the parent metal. As a result the weld may be anodic vs. the parent metal and may corrode more rapidly. This corrosion may take the form of localized metal wastage if H 2 S is present there is SSC cracking of hard zones in the metal or in the heat-affected zone. Similar problems can arise with electric- resistance-welded pipe resistance welded pipe. Metal wastage in sweet systems is avoided by using weld consumable with a higher alloy content than that of the base metal recourse is made to laboratory measurements to achieve the proper weld-metal/base-metal combination. Welding procedure standards are available to avoid hard zone combination. Welding procedure standards are available to avoid hard zone SSC. Chemical inhibition is also effective in protecting welded pipe. Fion Zhang/ Charlie Chong

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Preventing corrosion Th th t b iti i bl t ll t ihtf d The paths to obviating corrosion problems are conceptually straightforward:  Isolate the metal from the corrodant  Employ a metal alloy that is inherently resistant to corrosion in the py y y corrosive medium  Chemically inhibit the corrosion process  Move the electrical potential of the metal into a region where the corrosion  Move the electrical potential of the metal into a region where the corrosion rate is infinitesimally small “cathodic protection” An alternative is to live with the corrosion and replace the corroded component after failure after failure. Fion Zhang/ Charlie Chong

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Isolation I l ti i th i f i t ti d li A i t d ti t th Isolation is the regime of paints coatings and liners. An introduction to the subject is given in NACE 3 from which some of the following discussion is abstracted a detailed discussion of these subjects is in NACE. 10 For any j y coating to be effective it must be sufficiently thick to completely isolate the item being protected from the environment. Small holes in the coating “holidays” result in the rapid formation of pits Considerable care and quality holidays result in the rapid formation of pits. Considerable care and quality control is required to guarantee the generation of holidays during service. Organic coatings such as asphalt enamel and coal tar enamel are used to protect equipment concerned with the handling of oil and gas. Baked thin-film coatings such as thermosetting phenolics and epoxies applied in multiple coats can be used to protect tubular goods External protection of pipelines coats can be used to protect tubular goods. External protection of pipelines frequently involves use of adhesive tapes made of polyethylene or similar materials. Fusion bonded epoxy has been used successfully to protect a 150- km seawater-injection line oxygen was the corrodant much of which but not all was removed by scavenging chemicals. Fion Zhang/ Charlie Chong

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Inorganic coatings include both sacrificial coatings which furnish cathodic protection see below for mechanism at small breaks in the coating and protection see below for mechanism at small breaks in the coating and nonsacrificial coatings which protect only the substrates actually coated. Sacrificial coatings include galvanizing or coating with other metals anodic to the substrate and heavy suspensions of anodic metals e.g. zinc particles in silicates or organic vehicles. Zinc-silicate coatings paints are often used to coat the splash zone of drilling and production platforms The zinc metal coat the splash zone of drilling and production platforms. The zinc metal provides for cathodic protection of the steel substrate. Below the water line the most economical approach to corrosion control is cathodic protection see below. The pH of the environment is important— highly basic or acidic environments can remove coatings. Fion Zhang/ Charlie Chong

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Zinc-silicate coatings Fion Zhang/ Charlie Chong

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Nonsacrificial inorganic coatings include metal platings such as nickel and nonmetallic coatings such as ceramics Nickel can be applied by electroplating nonmetallic coatings such as ceramics. Nickel can be applied by electroplating or electroless plating. Ceramic coatings when properly applied are highly effective but they are also costly and fragile. Other systems while not truly coatings perform the same function e.g. Portland cement and plastic liners. Plastic liners have been used for internal protection of tubing and lined pipe. Some liners are sealed into individual joints of pipe and tubing some are Some liners are sealed into individual joints of pipe and tubing some are fused into one continuous close-fitting liner through the entire pipe. Both cement and plastic liners are suitable for water lines. The proper application of coatings is in large part an art form. Accordingly it is also not possible to overemphasize the need for close inspection of the coating process good quality control and testing that the coating has been coating process good quality control and testing that the coating has been complete. Fion Zhang/ Charlie Chong

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Pipeline Wrapping Fion Zhang/ Charlie Chong

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Corrosion resistant alloys CRA Ftitfil ll t l f d I t i From a cost point of view low-alloy steels are preferred. In certain cases “minor” alterations in alloy composition can minimize corrosion. For example L-80 steel with a tempered martensitic structure and a chromium content p 0.5 has been used without problems in 20-ppb oxygen-containing environments while a similar steel with 0.1 Cr has shown serious corrosion corrosion. The choice of using CRAs or chemical means to solve the more severe corrosion problem comes down to economics available capital vs. long-term operating costs. Remoteness of operation becomes an important consideration in determining operating costs as does downtime and consideration in determining operating costs as does downtime and deferred/lost oil because of repeated intervention for inhibitor application. Availability and cost of platform space is a consideration for offshore facilities. Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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The corrosion-control effectiveness of CRAs depends on the chemical severity of the environment Crevice corrosion pitting attack and SCC are the primary of the environment. Crevice corrosion pitting attack and SCC are the primary concerns. The corrosion resistance of annealed austenitic stainless steels such as 304 and 316 is affected by the presence of chlorides and temperature type 304 is less corrosion resistant than type 316. Both materials are susceptible to SCC when the temperature is above 150 °F. Both alloys are also low-strength steels Alloys 654 SMo and AL6XN can be manufactured to also low strength steels. Alloys 654 SMo and AL6XN can be manufactured to higher strengths and are more resistant to SCC. Austenitic stainless steels are probably the most susceptible of all ferrous alloys to pitting. Martensitic stainless steels have had the widest range of use of any of the available CRAs. Such steels may be manufactured through heat treatment available CRAs. Such steels may be manufactured through heat treatment into tubular products with acceptable yield strengths for downhole tubing. Many millions of feet of tubing type grade L-80 13Cr are in corrosive well iiti id dth tilfhi f d t ll ith service it is considered the material of choice for deep sweet-gas wells with temperatures less than 150 °C. About 35 of the L-80 13Cr usage was for oil wells. The passivity of 13Cr is destroyed by high chloride levels particularly at py y yg p y high temperature which can lead to pitting and crevice corrosion. Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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Duplex stainless steels are high-strength alloys achieved by means of cold working Such steels are more corrosion resistant than martensitic steels but working. Such steels are more corrosion resistant than martensitic steels but are similar in resistance to SSC. Cold-worked duplex has been used to 0.3 psi H 2 S. Annealed duplex is more resistant to H2S and SSC than the cold-worked versions. Annealed duplex line pipe has been used in wet CO 2 service 99 without problems. 22Cr duplex steel has been used where p H2S was between 0 5 and 1 psi Such steels have been used successfully in HT/HP wells e g 0.5 and 1 psi. Such steels have been used successfully in HT/HP wells e.g. 350 °F and 14000 psi producing no H 2 S. However the copresence of chloride stress and dissolved oxygen can induce SSC. Wells not exposed to even 13 small amounts of oxygen have operated successfully. 13 The material most commonly used for sour service is AISI Type 4130 steel The material most commonly used for sour service is AISI Type 4130 steel modified by microalloy additions with a quenched and tempered microstructure martensite. 14 C-110 steel has been used as casing in North Sll30t60bCO d 30 t 50 illib H S 15 Ai f Sea wells 30 to 60 bar CO 2 and 30 to 50 millibar H 2 S. 15 An overview of CRAs and their use in sour service is given in Treseder and Tuttle. Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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North Sea wells Fion Zhang/ Charlie Chong

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North Sea wells Fion Zhang/ Charlie Chong

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North Sea wells Fion Zhang/ Charlie Chong

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North Sea wells Fion Zhang/ Charlie Chong

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Nickel and cobalt alloys are used in the most severely corrosive conditions most severely corrosive conditions high pressure high temperature and high H 2 S contents. C-276 a nickel- based alloy can be used to 8000 psi H 2 S and 400 °F. Nickel alloys have found extensive use in the Mobile Bay found extensive use in the Mobile Bay fields. They are less expensive than the cobalt alloy MP35N previously used for such extreme conditions. Nickel alloys are also used as weld cladding for wellhead and valve equipment. wellhead and valve equipment. Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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Chemical inhibitors As with scale problems the appropriate addition of chemicals can often inhibit As with scale problems the appropriate addition of chemicals can often inhibit corrosion problems including some effects of H 2 S. The delivery techniques are often the same but the inhibition mechanisms and types of chemicals are different. Neutralizing inhibitors reduce the hydrogen ion in the environment. Typically they are: they are:  Amines  Ammonia  Morpholine They are effective in weak acid systems but are stoichiometric reactants: one They are effective in weak acid systems but are stoichiometric reactants: one molecule equivalent of inhibitor per molecule of acid. They have found minimal use in the oil field. Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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Scavenging inhibitors are compounds that also remove the corrodant. Oxygen scavengers are commonly used in the oil field e g in removing oxygen scavengers are commonly used in the oil field e.g. in removing oxygen during water injection. The majority of the corrosion inhibitors employed during production form thin barrier layers between the steel surface and the corroding fluid. The concept is that the organic inhibitor will strongly adsorb on the metal wall to form a barrier that the organic inhibitor will strongly adsorb on the metal wall to form a barrier possibly only a few molecules thick which will prevent access to the corrodant and possibly leave the surface oil-wet further retarding access of the corrodant. The generic name given to these compounds is “filming amines.” This name is qualitatively correct in that most inhibitors are indeed nitrogen- containing and the inhibitor does finally reside on the surface. The specific containing and the inhibitor does finally reside on the surface. The specific mechanism can be more complicated. For example the inhibitor can interact with the corrosion product to increase its adherence and to lower its bilit S h l lik l t b f thi k th f l l 7 permeability. Such layers are likely to be far thicker than a few molecules. 7 Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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Regardless of the specific mechanisms involved the inhibitor must contact the metal substrate The general procedures are: metal substrate. The general procedures are:  Tubing displacement  Displacement from the annulus  Continuous injection  Squeeze into the reservoir as liquid or gas  Weighed liquids/capsules/sticks  Weighed liquids/capsules/sticks  Vapor-phase inhibitors Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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The first two batch treatments are operated by pushing the inhibitor-containing fluid across the face of the production tubulars top down Item 1 or bottom up fluid across the face of the production tubulars top-down Item 1 or bottom-up Item 2. The inhibitor film then persists on the metal surface for some period of time ranging from days to months depending on the specific environment and materials. Continuous injection is done if the well completion allows for a “macaroni Continuous injection is done if the well completion allows for a macaroni string” reaching to the perforations. This technique often includes a simple-to- complicated valving system it should be remembered that valves can plug. Injection through the annulus has also been used. Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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Inhibitor squeezing into the formation is an alternative. Here the mechanism is different than that of scale inhibitor squeezes The large amount of inhibitor different than that of scale inhibitor squeezes. The large amount of inhibitor that returns initially is not wasted but is intended to coat the tubular and production equipment with an adsorbed persistent film of inhibitor. The small amounts of inhibitor that subsequently desorb from the formation are intended to repair holes that are generated in the initial film. Weighed liquids/capsules/sticks are all variations on the theme of placing inhibitor in the rathole where it is slowly released into the wellbore fluid continuously depositing and/or repairing the protective film. Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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Vapor-phase corrosion inhibitors are organic compounds that have a high vapor pressure generating volatile corrosion inhibitors such as some amines vapor pressure generating volatile corrosion inhibitors such as some amines that allow this inhibitor material to migrate to distant and often otherwise inaccessible metal surfaces within the container. Such inhibitors have been used on the Trans-Alaska pipeline to protect low-flow areas dead legs and the annular space in road casings and contingency equipment. The concept has also been applied to storage tank protection 16 has also been applied to storage tank protection. Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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Filming-amine inhibitors are intended to protect steels from the action of “natural” corrodants in the produced hydrocarbon and water phases They are natural corrodants in the produced hydrocarbon and water phases. They are generally not effective in protecting the steels from the acids used to stimulate wells or from the partially spent acids returning from such treatments. These tasks are accomplished by the inclusion of large dosages of different inhibiting chemicals with the stimulation acids. Such inhibitor systems are also available to handle low-alloy steels and CRAs in HT/HP conditions 17 Concern for to handle low alloy steels and CRAs in HT/HP conditions. Concern for stability of CRAs during matrix stimulation of deep hot wells has resulted in the use of organic acids such as acetic acid and formic acid rather than HCl. Inhibitor systems have been developed for these chemicals as well. Note: Matrix Stimulation http://www.glossary.oilfield.slb.com/Terms/m/matrix_stimulation.aspx A treatment designed to treat the near wellbore reservoir formation rather than other areas of the A treatment designed to treat the near-wellbore reservoir formation rather than other areas of the production conduit such as the casing across the production interval production tubulars or the perforations. Matrix stimulation treatments include acid solvent and chemical treatments to improve the permeability of the near-wellbore formation enhancing the productivity of a well improve the permeability of the near-wellbore formation enhancing the productivity of a well. Matrix stimulation is a process of injecting a fluid into the formation either an acid or solvent at pressures below the fracturing pressure to improve the production or injection flow capacity of a well. well. Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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Cathodic protection Thi t h l i d t t t i li ff h l tf d f This technology is used to protect pipelines offshore platforms and surface equipment. Corrosion is an electrochemical process:  Iron atoms give up electrons gp  Electrons flow through the metal to the corrodant  Ion movement in the water film contacting both corrodant and iron metal completes the electrical circuit completes the electrical circuit Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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In certain important cases it is possible to reverse this current flow out of the steel surface by the application of an external power supply i e make the steel surface by the application of an external power supply i.e. make the surface to be protected cathodic rather than anodic. The technology involved in employing cathodic protection must take into account:  Quantity of current required  Composition and configuration of the impressed current anode  Resistivity of the corroding medium  Resistivity of the corroding medium  Size of the item being protected  Accessibility of the surface being protected  Length of the item being protected Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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Fion Zhang/ Charlie Chong http://petrowiki.org/Corrosion_problems_in_production

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Fion Zhang/ Charlie Chong Fion Zhang Xitang 2016

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Fion Zhang/ Charlie Chong

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Read More Fion Zhang/ Charlie Chong

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1 SULFIDE STRESS CRACKING – PRACTICAL APPLICATION TO THE OIL AND GAS INDUSTRY Becky L. Ogden Southwest Petroleum Short Course Texas Tech University 2005 ABSTRACT The phenomenon of sulfide stress cracking SSC can result in catastrophic failures of pressurized equipment and piping resulting in extensive damage injuries and possible fatalities. Sulfide stress cracking was first identified as a serious problem in the oil industry in the late 1950’s with the development of deeper sour reservoirs. The high strength materials required for these wells began to fail as a result of brittle fracture that was later identified as SSC. Research began on this phenomenon and a task group was formed which later became associated with the National Association of Corrosion Engineers NACE now known as NACE International. The T-1B committee of NACE published a recommended practice addressing the metallic material requirements for protection against SSC. This recommended practice was later issued in 1975 as the Materials Requirement MR-0175 known today as “Metals for Sulfide Stress Cracking and Stress Corrosion Cracking Resistance in Sour Oilfield Environments”. Recently NACE International has issued the International Standard MR0175/ISO 15156 addressing multiple forms of cracking associated with the presence of aqueous hydrogen sulfide. This paper will concentrate on and identify the requirements for SSC to occur and give designers and operators practical options for the prevention of SSC in equipment operating in an aqueous H 2 S environment. While this paper will primarily discuss SSC some insight will be given to address the concerns of other forms of cracking. Key Words: hydrogen sulfide cracking sulfide stress cracking SSC partial pressure heat affected zone post weld heat treatment hardness sour environment metallurgy INTRODUCTION Aqueous hydrogen sulfide H 2 S in oil and gas production operations can result in many challenges. H 2 S is a poisonous gas that can result in severe metal loss corrosion as well as catastrophic brittle fractures of pressurized equipment and piping. These brittle fractures to metallic structures can happen quickly with little to no warning or may take years of exposure to occur. Several variables can influence a material’s likelihood or its resistance to cracking from exposure to hydrogen sulfide. The physical properties of the material the chemical properties of the material and the environment to which it is exposed all play an important role in determining whether a material is susceptible to SSC. Sulfide stress cracking or SSC is defined by NACE as the “Cracking of a metal under the combined action of tensile stress and corrosion in the presence of water and H 2 S a form of hydrogen stress cracking.” Through the review of this definition several factors must be present for SSC to occur. These factors are 1 a susceptible material 2 tensile stress 3 hydrogen sulfide and 4 water. If any one of these factors is missing sulfide stress cracking will not occur.

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2 MATERIAL PROPERTIES Steel is essentially a combination of iron and carbon with minor amounts of alloying elements added that enable that iron/carbon combination to perform the mechanical and chemical requirements of a particular grade of steel or alloy. The primary elements added include manganese silicon phosphorus chromium nickel and molybdenum. Each of these elements is added in varying concentrations so as to enhance the steel’s properties. However with regards to sulfide stress cracking or other forms of cracking these alloying elements must be reviewed and in some cases minimized. Materials have to be strong enough to perform under the conditions we require for our production conditions and designs. However generally with strength comes brittleness. Steels must be strong to perform yet ductile enough to prevent brittle fractures. A delicate balance must be obtained. As a result of laboratory testing and field experience NACE MR-0175:2003 details the parameters of acceptable chemical composition physical properties manufacturing processes and fabrication processes that will yield a material acceptable for use in a NACE defined sour environment. These parameters as they pertain to carbon steel materials will be detailed in a later section of this paper. FACTORS AFFECTING SULFIDE STRESS CRACKING Generally speaking an environment that produces hydrogen sulfide is considered “sour”. However for the environment to be defined as NACE sour it must exhibit characteristics that are favorable for the initiation of sulfide stress cracking. NACE MR0175:2003 defines the conditions in which SSC can occur. For the purpose of this discussion a “sour” environment shall be one in which the conditions are conducive to cracking by hydrogen sulfide. It is important to understand that the environments which can result in the SSC of materials are very specific in their compositions. SSC does not occur under all operating conditions. Several factors affect whether or not SSC will occur to a particular metallic structure. These factors include the alloy composition the material’s yield strength and hardness properties heat treatment microstructure fluid pH partial pressure of H 2 S total applied tensile stress and cold work temperature and time. How each of these impacts the SSC potential is discussed below. Alloy Composition The composition of a metallic material determines its susceptibility or resistance to various forms of cracking when exposed to particular environments. Generally speaking iron based materials or ferrous metals are more susceptible to SSC than nickel based alloys or non-ferrous materials. Additionally various levels of resistance/susceptibility to SSC can be found within a given family of alloys due to chemical compositional differences. Therefore each material should be reviewed prior to use to ensure it is acceptable for the intended use. Yield Strength and Hardness Properties In general the higher the strength of an alloy the harder the material and the more susceptible it is to sulfide stress cracking. Although yield strength is a true material property hardness is not. However

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3 there is a correlation between the two measurements. Generally the higher the yield strength the higher the hardness value. In most commercial grades of ferrous alloys the maximum strength level suitable for sour service use is 90000 psi yield strength. This roughly correlates to 22 Rockwell C HRC or 235 Brinell BHN hardness. This is often quoted for ferrous steels used in a NACE sour service. However through controlling steel chemistry and using special mill processing this upper limit can be increased. Testing and qualification of materials can be performed to determine its suitability for use in sour systems. Although hardness is not a true material property it is the preferred method of testing because it is simple and easy to perform relatively non-destructive and in most cases portable. Hardness values can be utilized by manufacturers and procurement agents as a quality control method during the fabrication process or by the field personnel as a field inspection technique. Additional discussions on hardness determinations are discussed in the next section. Heat Treatment The type of heat treatment applied to a particular alloy can affect the material’s microstructure and ultimately its susceptibility to sulfide stress cracking. A microstructure comprised of tempered martensite with fine grains will result in materials of superior resistance to SSC. Carbon and low alloy steels are acceptable in the as-milled condition as long as they contain less than 1 nickel meet the hardness requirements and are in one of the following heat-treatment conditions: hot-rolled carbon steels only annealed normalized normalized and tempered normalized austenitized quenched and tempered or austenitized quenched and tempered. It should be noted that field fabrication cold working and welding of “approved” materials can alter the microstructure making the material susceptible to SSC. It may be necessary to thermally stress relieve the materials following these processes to “reinstate” their resistance to SSC. Microstructure Although susceptibility to SSC increases with increasing hardness some microstructures are more susceptible to cracking than others at the same hardness levels. As stated above the tempered martensite is more resistant to SSC than the tempered bainite or mixed structures of the same hardness. Additionally the degree of segregation and the type size shape and distribution of inclusions are other microstructural variables that can influence the resistance to sulfide stress cracking. Fluid pH The higher the fluid pH the more resistant materials are to SSC. This tendency enables drilling operations to utilize high strength materials in zones known to produce H 2 S. Although pH control is acceptable and manageable in drilling operations it is not readily utilized in production scenarios. Maintaining a constant pH in production would prove troublesome and impractical. Therefore hardness limitations and alloy selections are the preferred method for controlling SSC.

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4 Partial Pressure of H 2 S As the partial pressure of H 2 S increases the susceptibility of a material to SSC increases. The partial pressure of H 2 S is defined as the portion of the total pressure associated with the specific component of interest in this case H 2 S. The partial pressure is calculated by multiplying the total system pressure by the mole fraction of H 2 S in the gas phase. If the calculated partial pressure of H 2 S is above 0.05 psia in a gas system SSC is possible. Figures 1 and 2 show the relationship to H 2 S system pressure and partial pressure for gas and multiphase systems as illustrated by NACE MR0175:2003 edition. It should be noted that these limits are a “practical” limit due to other factors affecting SSC materials have failed at partial pressures below 0.05 psia. Therefore care should be taken to review all factors involved in the material selection. Total Applied Tensile Stress and Cold Work Different alloys possess different threshold levels at which SSC will occur. Understanding this threshold level will enable the designer to ensure that the stresses applied to a material will not result in cracking. It needs to be understood that the total stresses working on a material are the combination of both the applied stress i.e. pressure and the residual stress fabrication/manufacturing stresses. The higher the applied stress on a material the more susceptible to SSC it becomes. Cold work or cold formed materials may be susceptible to SSC at hardness levels below HRC 22 BHN 235. Cold working will alter the microstructure and increase the residual surface tensile stresses. For this reason heat treatment is recommended for cold worked or cold formed low alloy steels before they are used in a sour environment. An annealing or normalizing heat treatment will return the material to its original SSC resistance following cold working. Temperature The potential for SSC decreases as temperatures increase. Therefore additional high strength tubing and casing materials can be utilized above threshold temperatures. However if a well is to be completed or operated in a sour zone with a temperature above a threshold temperature for a particular material the engineer must confirm that the environment in contact with the material does not drop below that critical temperature. Below this temperature these high strength materials are susceptible to SSC and cannot be utilized. Table 1 an excerpt from the NACE MR0175:2003 standard illustrates the temperature dependence of tubing and casing materials in oil and gas wells. Time The general rule of thumb is “the longer the time of exposure at a constant stress level the greater the danger of SSC for susceptible alloys”. Under laboratory controlled conditions it is possible to determine the time to failure of a given alloy under a particular set of conditions. However in actual field conditions projecting a time to failure is extremely difficult. The time it takes for a material to fail due to SSC is dependent on the aggressiveness of the environment and the degree of susceptibility of the material. SSC can happen quickly or may take years to develop. Therefore it is critical that a review of the materials and environment be conducted prior to specifying the completion equipment. SSC resistant materials should be utilized.

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5 HARDNESS TESTING OF MATERIALS While hardness testing is a simple procedure it must be performed correctly and must represent the material in the as-received or as-fabricated condition. Hardness by definition is the resistance of a metal to plastic deformation usually by an indention. Hardness testers utilize an indenter which is forced into the metal surface by a known loading. The relationship to the area or depth of the indentions to the load applied is known as the hardness of the material. Hardness can be measured on multiple “scales”. NACE MR0175 utilizes the Rockwell C scale HRC or the Brinell scale BHN. If hardness values are specified for the parent metal and any heat affected zones left in the as-welded or as- milled condition there shall be a sufficient number of hardness tests performed to ensure the readings are below the specified value as noted within NACE MR0175/ISO 15156:2003 for that particular material. Controlling hardness is an acceptable method for preventing SSC. MR0175/ISO 15156:2003 does not specify the number or locations of hardness tests on the parent material. However if hardness control is to be utilized for approving a welding procedure for use in sour services specific locations and numbers of tests must be performed. These are noted within the International NACE MR0175/ISO 15156:2003 standard for Vickers and Rockwell Hardness measurements for fillet welds butt welds and repair/partial penetration welds. These illustrations must be followed for weld procedure qualifications. Figures 3 and 4 illustrate the survey method requirements for hardness measurements on butt welds. WELDING AND ITS IMPACT ON SSC Welding is a “necessary evil” in sour systems in the oil and gas industry however steps can be taken to minimize its negative impacts. When steel is welded the parent material and consumables are variables that must be reviewed and controlled. But these are not the only variables that need to be considered when welding in sour service. The effects of rapid cooling in the heat affected zone HAZ of a weld can result in areas of localized hardness. The HAZ is that area around the actual weldment that has been exposed to high temperatures but not high enough to actually liquefy the material. However as a result of this heat input phase transformations do occur resulting in a microstructure has been “partially melted” and altered due to the heat of welding. This altered HAZ is now more susceptible to SSC due to its increased hardness. A parent material that was suitable and acceptable in regard to SSC in its as-received as-milled condition may be susceptible to SSC following fabrication that involves welding. Therefore fabrication processes involving welding must be reviewed for their potential impact on the SSC potential of the parent material. Weld procedures can be written and qualified as being in compliance with NACE with regard to both SSC and other forms of cracking associated with the presence of aqueous hydrogen sulfide. The qualification requirements for hardness measurements traverse across the weld are detailed in the NACE MR0175/ISO 15156:2003 Standard as stated in the previous section on hardness measurements again see Figures 3 and 4. However in lieu of qualifying a weld procedure to NACE one also has the option to post weld heat treat PWHT following the completion of the welding. The use of a PWHT technique tempers the welded material. This reduces the residual internal stresses created when the weld metal solidified and tempers any martensite that may be present into a configuration of lower internal strain. The process of

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6 PWHT will be specific for each type and thickness of material and the procedures are described within ASME Boiler and Pressure Vessel Code Section 8 Division 1. When specifying line pipe for sour service it has been this author’s experience to prefer the seamless line pipe over Electric Resistance Weld ERW pipe. When ERW pipe was utilized the specification always called for full body normalizing following manufacturing verses only seam annealing following manufacturing. In my experience this proved to provide better resistance to SSC when exposed to the severe H 2 S environments in the Permian Basin area of West Texas and Southeastern New Mexico. USE OF PLATINGS AND COATINGS While platings and coatings are an acceptable barrier for generalized corrosion they are not acceptable for use in the prevention of SSC as per NACE MR0175-2003. DETERMINING A SOUR ENVIRONMENT Hydrogen sulfide is one of the most serious corrosion agents encountered in the oil and gas industry. In addition to its ability to crack metals it can also result in pitting corrosion with subsequent failures. The release of H 2 S as a result of corrosion or cracking can endanger the lives of people working around or in near proximity to the release point. H 2 S can be fatal at concentrations as low as 500 ppm. Therefore designing equipment resistant to H 2 S cracking is critical. Additionally prevention of corrosion by H 2 S is also highly recommended. Inhibition of corrosion can be obtained through material selection internal coatings or the application of corrosion inhibitor. However to prevent cracking the NACE standard must be strictly adhered to and followed. It is the responsibility of the owner/user to determine whether a given environment falls within the parameters of a sour environment thus requiring SSC resistant materials. Information concerning the environment’s operating pressure H 2 S content water content pH and temperature all play a role in making this determination. Designing for SSC resistance is not only a prudent and good engineering practice but it is a requirement of many regulatory agencies. Specifically the Texas Railroad Commission Rule 36 the BLM On-Shore Order 6 and the New Mexico Statewide Rule 118 all specify that SSC resistant materials must be utilized in an H 2 S environment. Therefore this is both critical for safety and regulatory compliance Note: Currently the regulations still specify NACE MR0175 latest edition as the standard for compliance. It is unknown as of the writing of this paper as to whether the agencies will adopt the new International MR0175/ISO 15156 Standard. However producers should be aware of the changes published in the new standard and be prepared to make appropriate modifications to fabrication and engineering specifications. Referring again to Figures 1 and 2 will enable the user/owner to evaluate his/her system based on H 2 S content and pressure assuming the presence of free water. It should be noted that hydrogen sulfide can be present naturally in produced fluids or can be introduced as a result of contamination by incompatible waters or sulfate reducing bacteria. Frequent surveys of non-sour or “sweet” fluids should be conducted to determine if hydrogen sulfide generation is

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7 occurring over the life of a well or producing field. All safety precautions should be exercised when determining the concentration of H 2 S in production fluids. Because of the dangers associated with low concentrations of H 2 S it is recommended to always assume H 2 S is present regardless of the past history of a field lease or individual well. OTHER FORMS OF HYDROGEN DAMAGE In addition to SSC there are other forms of hydrogen damage and cracking that can occur in aqueous hydrogen sulfide environments. Because the potential for catastrophic failures associated with these forms of cracking also exists NACE has recently published a joint international standard NACE MR0175/ISO 15156:2003. This standard addresses the concerns for all types of cracking associated with sour production and makes recommendations for materials and operating conditions to prevent such failures. Hydrogen Induced Cracking HIC Hydrogen Induced Cracking HIC is defined as a “hydrogen attack induced by decarburization”. This type of attack occurs at elevated temperatures and is caused by atomic hydrogen permeating through the steel and reacting to form other gases. Hydrogen reacts with the carbon in the steel to form methane gas which cannot diffuse out of the steel’s matrix. Accumulation of this methane at grain boundaries and other steel discontinuities results in localized high stresses from which cracks can occur. HIC attack generally occurs at temperatures greater than 500 0 F and is dependant on the hydrogen partial pressure. Consult the API Publication 941 Steels for Hydrogen Service at Elevated Temperatures and Pressures in Petroleum Refineries and Petrochemical Plants for additional information on this phenomenon. Step Wise Cracking SWC Step Wise Cracking SWC is defined as “hydrogen cracks which lie parallel to each other and are connected by cracks between them”. This type of cracking can be either discrete cracks or an array of cracks. The cracks that connect the “main cracks” and lead to SWC are caused by the shear stresses between the main cracks. This type of cracking can lead to catastrophic failures due to the potential for the cracks to propagate through the thickness of the material resulting in a considerable loss of strength and ultimate failure. Stress Oriented Hydrogen Induced Cracking SOHIC Stress Oriented Hydrogen Induced Cracking SOHIC is defined as “hydrogen induced cracking propagated by high internal stresses typically hoop stress”. This type of cracking is similar to HIC and SSC but the cracking is transgranular or across the grains in the through thickness direction. These cracks initiate and propagate in the direction normal to the applied stress and are typically observed in the HAZ of relatively high hardness microstructures. The application of a high external stress i.e. pressure typically contributes to the failure.

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8 Hydrogen Blistering Hydrogen Blistering is defined as the “subsurface cracking from absorption and concentration of hydrogen”. Blistering occurs when hydrogen enters the steel and combines into molecular hydrogen at defects present in the steel plate typically non-metallic inclusions such as sulfides. Hydrogen blistering generally occurs in low pressure equipment such as tanks and pipeline equipment that are exposed to a corrosive environment that contains hydrogen sulfide. Because of the nature of the manufacturing process of rolled plates inclusions present in the steels become elongated with the larger inclusions present and aligned along the centerline of the plate. It is at these larger inclusions that hydrogen blistering tends to occur. The high internal pressures present with the formation of the molecular hydrogen create high internal stresses within the steel that can greatly exceed the yield strength of a material and result in the formation of blisters. These blisters can often be visually observed on the exterior surface of steels in the form of an area of localized “swelling”. They often resemble a “paint blister” but they are indeed a blister in the steel plate. The best prevention for this type of cracking is to specify “HIC resistant material” when procuring steel plates. This material has substantially lower sulfur content usually 0.005 maximum sulfur and is commonly calcium-treated for sulfur shape control. This lower sulfur content reduces the amount of inclusions present in the steel thereby reducing the number of available sites for hydrogen to accumulate and form molecular hydrogen. The calcium treatments help to “round” any inclusions that may be present in the low sulfur steel thereby making it more difficult for hydrogen to enter. CONCLUSIONS It is the goal of design engineers and operators to prevent failures whether they are annoying seeps or catastrophic failures. It is a considerable economic benefit for users of steels to understand the environment in which that steel will be placed and the hazards associated with its exposure to that environment. By understanding the various forms of cracking that can occur to steels the designer/operator can implement specification changes or modify the environment to eliminate that mechanism. This will extend the effective life of the equipment reduce the potential for failures reduce downtime associated with equipment repairs and make the production area safer with respect to equipment failure incidents. It is vital when selecting a manufacturing/fabrication process to consider the preventative measures for reducing the susceptibility of steel to various forms of cracking. Some Regulatory agencies require that equipment be manufactured fabricated and maintained in a condition that is resistant to Sulfide Stress Cracking SSC. Therefore it is critical during the early stages of a project to specify “NACE” trim and “NACE” compliance on all equipment in a sour environment. However it must be noted that any modifications made to such equipment following its installation must be made such that this “NACE” condition is not negated. By following the Standards published by NACE the designer/operator can be confident that the specified equipment is acceptable for use in a sour environment and is resistant to SSC. By applying the details found in the new International Standard NACE MR0175/ISO 15156: 2003 additional protection from other forms of cracking can be integrated into the design specifications.

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28 Table 1: Acceptable API and ASTM Specifications for Tubular Goods Copyright NACE International 2003 Used with Permission

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29 Table 2 — List of equipment Copyright NACE International 2003 Used with Permission NACE MR0175/ISO 15156 is applicable to materials used for the following equipment Permitted exclusions Drilling well construction and well-servicing equipment Equipment only exposed to drilling fluids of controlled composition a Drill bits Blowout Preventer BOP shear blades b Drilling riser systems Work strings Wireline and wireline equipment c Surface and intermediate casing Wells including subsurface equipment gas lift equipment wellheads and christmas trees Sucker rod pumps and sucker rods d Electric submersible pumps Other artificial lift equipment Slips Flow-lines gathering lines field facilities and field processing plants Crude oil storage and handling facilities operating at a total absolute pressure below 045 MPa 65 psi Water-handling equipment Water-handling facilities operating at a total absolute pressure below 045 MPa 65 psi Natural gas treatment plants Transportation pipelines for liquids gases and multiphase fluids Lines handling gas prepared for general commercial and domestic use For all equipment above Components loaded only in compression a See A.2.3.2.3 for more information. b See A.2.3.2.1 for more information. c Wireline lubricators and lubricator connecting devices are not permitted exclusions. d For sucker rod pumps and sucker rods reference can be made to NACE MR0176. Key X H 2 S partial pressure kPa Y in situ pH 0 Region 0 1 SSC Region 1 2 SSC Region 2 3 SSC Region 3 In defining the severity of the H 2 S-containing environment the possibility of exposure to unbuffered condensed aqueous phases of low pH during upset operating conditions or downtime or to acids used for well stimulation and/or the backflow of stimulation acid after reaction should be considered.

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30 Figure 1: H2S Partial Pressure Relationship for a Gas System Copyright NACE International 2003 Used with Permission

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31 Figure 2: H2S Partial Pressure Relationship for a Multi-Phase System Copyright NACE International 2003 Used with Permission

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32 Figure 3: Butt weld survey method for Vickers hardness measurement Copyright NACE International 2003 Used with Permission

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33 Figure 4: Repair and Partial Penetration Welds Copyright NACE International 2003 Used with Permission

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34 List of References: NACE MR0175-2003 “Metals for Sulfide Stress Cracking and Stress Corrosion Cracking Resistance in Sour Oilfield Environments” NACE International Houston TX:NACE NACE MR0175/ISO 15156:2003 “Petroleum and Natural Gas Industries – Materials for use in H2S Containing Environments in oil and gas production” NACE International Houston TX:NACE McIntyre D. R. Moore E. M. Jr. : “Specified Pipe Fittings Susceptible to Sulfide Stress Cracking” Materials Performance January 1996 pp 64 – 66 Tsukano T. et al “Development of Sour Service Drillstring with 110-ksi Yield Strength” SPE/IADC Drilling Conference 1991 No. 22004 Tuttle R. N. “What Is a Sour Environment” Journal of Petroleum Technology March 1990 pp 260 - 262 Texas Administrative Code Title 16: Economic Regulation Part 1 Railroad Commission of Texas Chapter 3: Oil and Gas Division Rule 3.36 Oil Gas or Geothermal Resource Operation in Hydrogen Sulfide Areas 2004 Bureau of Land Management Onshore Oil and Gas Order No. 6 43 CFR 3160 Federal Register/ Volume 55 No. 226 New Mexico State Wide Rule 118: 19153.118 Hydrogen Sulfide Gas Hydrogen Sulfide

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Revision 3 PV Plates for sour Service Verfasser/Dokument 1 Requirements for steel plates in sour service

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Revision 3 PV Plates for sour Service Verfasser/Dokument 2 Sour Service an introduction in the world of hydrogen induced corrosion

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3 Requirements for steel plates in sour service Verfasser/Dokument 3 Sour service damage is not a new issue The oldest reports about sour service steel damage more than 60 years old Many organisations like NACE or EFC oil- and gas companies engineering companies are still improving regulations The importance of hydrogen damage due to sour service is more and more recognised. The exploitation of sour gases and out of sour oil sources is rising. Often sweet sources get more and more sour.

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4 Requirements for steel plates in sour service Verfasser/Dokument 4 Why this sensitivity to sour service damage Sour media are aggressive to steel structures damages not easy to detect. Health and safety of personnel and the public are in danger if precautions in survey of equipment and a right material selection are not adjusted. Severe environmental pollution could be the consequence out of such damages. Shutdowns due to material failures and the replacement of pressure vessels can cause dramatic economical loss. A really bad example: Accident at Chicago refinery in 1984 17 people killed. Many good reasons for our full attention

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5 Requirements for steel plates in sour service Verfasser/Dokument 5 Union Oil absorber vessel failure resulting from cracks growing in HAZ with no PWHT

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6 Requirements for steel plates in sour service Verfasser/Dokument 6 The view of the steel plate manufacturer Steel plate requisitions reflect an increasing demand for plates with improved properties for sour service Large variety of customer requests: - many specifications based on published recommendations or test methods e.g. NACE MR 0175 TM0284... - in combination with the “in house”-experience and -prescriptions Aim of this paper: - general overview over the damaging mechanisms - general survey about the current specified requisitions for plate orders - Dillinger Hütte GTS possibilities to supply improved steel plates

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Revision 3 PV Plates for sour Service Verfasser/Dokument 7 Damaging Mechanisms and Test Methods

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8 Requirements for steel plates in sour service Verfasser/Dokument 8 What are the sour service corrosion mechanisms Hydrogen-Induced Cracking HIC Hydrogen Blistering Sulfide Stress Cracking SSC probably to be taken into consideration: Stress-Oriented Hydrogen-Induced Cracking SOHIC

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9 Requirements for steel plates in sour service Verfasser/Dokument 9 Electrons Molecular Hydrogen Sulfide Ionics Proton Hydrogen Atom Steel with typical small imperfections Hydrogen Sulfide Acidic H 2 S -containing medium Cracking mechanism in the steel during H2S corrosion process

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1 0 Requirements for steel plates in sour service Verfasser/Dokument 10 H 2 S → 2 H + + S 2- Fe + 2 H + → Fe 2+ + 2 H ad Fe 2+ + S 2- → FeS H 2 S+ Fe → FeS + 2 H ab 2 H ab → H 2 Corrosion reaction

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1 1 Requirements for steel plates in sour service Verfasser/Dokument 11 HIC / SWC Blistering SSC SOHIC Schematical appearance of damage mechanisms in sour service

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1 2 Requirements for steel plates in sour service Verfasser/Dokument 12 HIC Hydrogen Induced Cracking Corrosion at stress free prismatic specimens Definition as per NACE MR0175/ISO 15156: Planar cracking that occurs in carbon and low alloy steels when atomic hydrogen diffuses into the steel and then combines to form molecular hydrogen at trap sites.

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1 3 Requirements for steel plates in sour service Verfasser/Dokument 13 “Evaluation of Pipeline and Pressure Vessel Steels for Resistance to Hydrogen-Induced Cracking” HIC: Stepwise internal cracking on different planes of the metal no external stress origin: 1984 for evaluation and comparison of test result test solution: pH 3 sol. A and pH 5 sol. B saturated with H 2 S test specimens: position one end/mid width preparation dimensions duration: 96 h evaluation: metallographic examination of cross sections acceptance crit.: to be agreed between purchaser and supplier documentation: CLR CTR CSR values for each section specimen test NACE TM 0284-2003

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1 4 Requirements for steel plates in sour service Verfasser/Dokument 14 Test specimen location acc. to NACE TM 0284

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1 5 Requirements for steel plates in sour service Verfasser/Dokument 15 test duration: 96h test solution: saturated with H 2 S test solution H 2 S 100 20 Test specimens HIC test method acc. to NACE TM 0284 Solution A -pH 3 - 5 NaCl 0.5 CH 3 COOH - identical to Solution A of NACE TM 0177 Solution B -pH 5 - synthetic seawater acc. ASTM D1141

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1 6 Requirements for steel plates in sour service Verfasser/Dokument 16 In Detail

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1 7 Requirements for steel plates in sour service Verfasser/Dokument 17 HIC test vessel test specimens during HIC-test

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1 8 Requirements for steel plates in sour service Verfasser/Dokument 18 sectioning of test specimens 20 mm 25 mm 25 mm 25 mm 25 mm rolling direction faces to be examined Examination of the polished sections: a W b b a a crack length b crack width W specimen length T T specimen thickness Crack distance 0.5 mm single crack 100 ⋅ ∑ W a CLR 100 ⋅ ∑ T b CTR 100 ⋅ ⋅ ⋅ ∑ T W b a CSR

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1 9 Requirements for steel plates in sour service Verfasser/Dokument 19 HIC or SWC damage

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2 0 Requirements for steel plates in sour service Verfasser/Dokument 20 A516 GR70 Amine Contactor 1 1: NACE RP0296 Hydrogen Blistering

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2 1 Requirements for steel plates in sour service Verfasser/Dokument 21 A516 GR70 Amine Contactor 1 1: NACE RP0296 Hydrogen Blistering

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2 2 Requirements for steel plates in sour service Verfasser/Dokument 22 Amine Contactor/Water Wash Tower 1 1: NACE RP0296 Blister Cracking

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2 3 Requirements for steel plates in sour service Verfasser/Dokument 23 SSC Sulfide Stress Cracking Corrosion at specimens under stress Definition as per NACE MR0175/ISO15156: Cracking of metal involving corrosion and tensile stress residual and/or applied in the presence of water and H 2 S

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2 4 Requirements for steel plates in sour service Verfasser/Dokument 24 „Laboratory Testing of Metals for Resistance to Specific Forms of Environmental Cracking in H 2 S Environments” origin: 1977 revised 1986 1990 and 1996 4 test methods: tensile test sol.A preferred by DH-GTS 1 Bent-Beam Test sol. B C-Ring test sol. A Double-Cantilever-Beam test DCB sol.A 2 test solutions: A: pH: 2.7 B: pH: 3.5 H 2 S saturated test duration: 720 h or until failure whichever occurs first results report: applied stress over log time stress level of no fail. after 720h remark DH-GTS: acceptable only if PWHT plus DICREST route no microalloying elements 1: also 4 point bend test acc. ASTM G39 sol.A typ. linepipe NACE TM0177

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2 5 Requirements for steel plates in sour service Verfasser/Dokument 25 Sulfide Stress Cracking SSC in HAZ of head to shell weld of FCC absorber tower.

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2 6 Requirements for steel plates in sour service Verfasser/Dokument 26 Sulfide Stress Cracking

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2 7 Requirements for steel plates in sour service Verfasser/Dokument 27 SSC four-point bend test

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2 8 Requirements for steel plates in sour service Verfasser/Dokument 28 SSC tensile test

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2 9 Requirements for steel plates in sour service Verfasser/Dokument 29 SOHIC Stress Orientated Stress Cracking Corrosion at notched specimens under stress Definition as per NACE MR0175/ISO15156: Staggered small cracks formed approximately perpendicular to the principle stress residual or applied resulting in a „ladder-like“ crack array linking sometimes small pre-existing HIC cracks.

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3 0 Requirements for steel plates in sour service Verfasser/Dokument 30 Sporadic documentation at spiral welded pipes and flaws in pressure vessels. Combination of rectangular SSC type and parallel cracks HIC type in the area of a multi dimensional tension field. Typical SOHIC crack below a flaw. Created in a double beam bend test. New phenomenon in the field of sour gas corrosion Stress-Oriented Hydrogen Induced Cracking SOHIC

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3 1 Requirements for steel plates in sour service Verfasser/Dokument 31 SOHIC-Crack at a non PWHT repair weld of a primary absorber deethanizer 1. 1: NACE RP0296 Stress-Oriented Hydrogen Induced Cracking SOHIC

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3 2 Requirements for steel plates in sour service Verfasser/Dokument 32 Stress-Oriented Hydrogen Induced Cracking SOHIC •issue - still under large discussion •mechanism not fully understood •mixture of SSC and HIC type cracking •location close to the welds

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3 3 Requirements for steel plates in sour service Verfasser/Dokument 33 SOHIC test as per NACE TM0103 / 2003 • SOHIC testing • 4 point bent double beam tests • test duration 168 h • metallographic examination of the cross sections • Reasonable acceptance criteria for CCL Continuous Crack Length DCL Discontinuous Crack Length and TCL Total Crack Length are not yet reported

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3 4 Requirements for steel plates in sour service Verfasser/Dokument 34 NACE TM0103 – Full Size Double-Beam Test Specimen Design SOHIC test arrangement as per NACE TM0103 / 2003

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3 5 Requirements for steel plates in sour service Verfasser/Dokument 35 Dimensions of the notch: Depth 2mm r 0.13mm Sectioning across the notch into two cross sections. cut line notch centred 5 cm Section 1 Section 2 Drop faces to be examined SOHIC test specimens as per NACE TM0103 / 2003

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3 6 Requirements for steel plates in sour service Verfasser/Dokument 36 CCL - continuous cracks perpendicular in the most stressed area near to the bottom of the notch. DCL - discontinuous parallel cracks below the continuous crack area with lower stresses. TCL - length of the whole cracked area. SOHIC evalutation of the cross sections from the double beam specimens

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3 7 Requirements for steel plates in sour service Verfasser/Dokument 37 Results of the SOHIC tests at Dillinger Hütte GTS 1 Although the tests were performed with HIC resistant DICREST material at a load of less than 50 yield in pH3 solution first SOHIC type cracks appeared. Rising the load increases the appearance of these cracks Testing in pH5 solution no SOHIC cracks are detected. The notch of specimens generates a very too harmful stressed area.

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3 8 Requirements for steel plates in sour service Verfasser/Dokument 38 It should be taken into consideration whether a notch like this is permitted generally at pressure vessels. This could explain why even HIC and SSC resistant steels DICREST show big amounts of SOHIC cracks with the proposed test method. Acc. to DH’s opinion this test method is not appropriate as SOHIC test. SOHIC resistant material acc. to this test method can not be produced with normalised steels. It seems to be that Q+T material will reach this aim. Results of the SOHIC tests at Dillinger Hütte GTS 2

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3 9 Requirements for steel plates in sour service Verfasser/Dokument 39 SSC + HIC Standards

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4 0 Requirements for steel plates in sour service Verfasser/Dokument 40 NACE MR0175/ISO 15156 - 2003 “Petroleum and natural gas industries—Materials for use in H 2 S- containing environments in oil and gas production” • By the end of 2003 NACE0175/ISO15156 was published giving requirements and recommendations for the selection and qualification of carbon and low- alloy steels corrosion-resistant alloys and other alloys for service in equipment used in oil and natural gas production and natural gas treatment plants in H 2 S- containing environments • 3 parts: - Part 1: General principles for selection of cracking-resistant materials - Part 2: Cracking-resistant carbon and low alloy steels and the use of cast irons - Part 3: Cracking-resistant CRAs corrosion-resistant alloys and other alloys • Qualification route for steels not yet proved to be suitable for H 2 S service

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4 1 Requirements for steel plates in sour service Verfasser/Dokument 41 SSC: Metal cracking under corrosion in presence of H 2 S and stress same time hydrogen embrittlement especially in steel with high hardness or high strength SSC and SCC susceptibility depends on e. g.: - steel: chemical composition heat treatment microstructure cold deformation - hydrogen activity pH-value - total tensile stress including residual stress - temperature duration ... Definition of SSC severity levels from 0 to 3 with increasing severity severity level 1starting from H 2 S partial pressure ≥ 0.0003 MPa No absolute resistance material can fail in SSC-tests SSC in NACE MR0175/ISO 15156 - 2003

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4 2 Requirements for steel plates in sour service Verfasser/Dokument 42 Requirements: Pressure vessel steels classified as P-No 1 group 1 or 2 in Section IX of the ASME Boiler and Pressure Vessel Code are acceptable without testing Carbon low alloy steels: - heat treated contr. Rolled N N+T Q+T - Ni 1 wt - Hardness 22 HRC average 24 HRC individual fabrication conditions: welding and PWHT have to respect 22HRC limitation also in HAZ and WM 5 cold deformation SR to be applied Remark of the steel producer: NACE MR0175 shall prevent SSC-Cracking but there is very few influence on steel making practice no influence on HIC-resistance SSC in NACE MR0175/ISO 15156 – 2003

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4 3 Requirements for steel plates in sour service Verfasser/Dokument 43 Spec Grade UNS Spec Grade UNS SA-283 A B C D - SA-299 ... K02803 SA-285 C K02801 SA-455 ... K03300 SA-285 A K01700 SA-515 70 K03101 SA-285 B K02200 SA-516 70 K02700 SA-36 ... K02600 SA-537 Cl. 1 K12437 SA-515 65 K02800 SA-662 C K02007 SA-515 60 K02401 SA-737 B K12001 SA-516 55 K01800 SA-738 A K12447 SA-516 60 K02100 SA-516 65 K02403 SA-562 ... K11224 SA-662 A K01701 SA-662 B K02203 P-No.1 Group 2 P-No.1 Group 1 Listing of Section IX of the ASME Boiler Pressure Vessel Code

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4 4 Requirements for steel plates in sour service Verfasser/Dokument 44 - The user shall consider HIC and HIC testing even if there are only trace amounts of H 2 S present - HIC susceptibility is influenced by chemistry and manufacturing route Requirements - low Sulphur content 0003 - test acc. to NACE TM0284 - acceptance criteria solution A: CLR ≤ 15 CTR ≤ 5 CSR ≤ 2 - other conditions may be defined as per table B.3 for specific or less severe duty HIC in NACE MR0175/ISO 15156 – 2003

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4 5 Requirements for steel plates in sour service Verfasser/Dokument 45 SOHIC in NACE MR0175/ISO 15156 – 2003 User should consider SOHIC when evaluating carbon steels - Pre-qualification to SSC prior to SOHIC/SZC evaluation - Small-scale tests: unfailed uniaxial tensile UT four point bend FPB specimen are metallographicly examined - UT-specimen : - no ladderlike HIC indications or cracks exceeding 05mm in through thickness direction allowed - after hydrogen effusion the tensile strength shall not be less than 80 of the tensile strength of unused specimens - FPB-specimen: - no ladderlike HIC indications or cracks exceeding 05mm in through thickness direction allowed - blisters less than 1mm below the surface and blisters due to compression regardless of the depth shall be disregarded - Full pipe ring tests may be used test method and acceptance criteria described in HSE OTI-95-635

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4 6 Requirements for steel plates in sour service Verfasser/Dokument 46 EFC 16 “Guidelines on Materials Requirements for Carbon and Low alloy Steels for H2S-Containing Environments in Oil and Gas Production Combined specification for test methods of HIC and SSC” concerns: C- and low alloy steels in oil and gas production not in refinery service conclusion of NACE-test methods published: in 1995 rev. 2 in 2002 1.HIC - low S shape control low segregation low CEQ - test acc. to NACE TM0284 Solution A - acceptance criteria: CLR ≤ 15 CTR ≤ 5 CSR ≤ 1.5

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4 7 Requirements for steel plates in sour service Verfasser/Dokument 47 EFC 16 2 2. SSC - fpH-value/ H2S-p.pressure: Non sour transition region sour service - in case of sour service: see guidelines limited hardness in HAZ to max. 250 HV30 except cap´s cap layer up to 275 HV30 t 95mm or 300 HV30 t 95mm limited cold deformation 5 for PV or PWHT 620°/650°C - various test methods for the evaluation of SSC resistance uniaxial 4point- bend C-ring.... pH 3 DH recommend the tensile test and 4 point bend test - load and duration of the test to be agreed proposals are made recommendation of DH-GTS e.g.: load: 0.72 SMYS duration: 720 h 3. SOHIC/ SZC Soft zone cracking - PWHT recommended - testing the susceptibility by 4 point bend test as an option however no acceptance criteria defined

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4 8 Requirements for steel plates in sour service Verfasser/Dokument 48 HIC + SSC Guidelines RP + MR

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4 9 Requirements for steel plates in sour service Verfasser/Dokument 49 “Guidelines for Detection Repair and Mitigation of Cracking of Existing Petroleum Refinery Pressure Vessels in Wet H 2 S Environments” Concerns HIC SSC SOHIC ASCC Alkaline Stress C.C. applicable for existing equipment in refineries made of carbon steel valid if H 2 S concentration ≥ 50 ppm but no threshold concentration defined reports about the parameters for each damage mechanism reports about a large survey in 1990 of 5000 inspected pressure vessels 26 of all vessels showed cracking incidence crack depth from 1.6 mm to more than 25 mm recommendations for inspection NACE RP0472 – 2000

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5 0 Requirements for steel plates in sour service Verfasser/Dokument 50 “Guidelines for Detection Repair and Mitigation of Cracking of Existing Petroleum Refinery Pressure Vessels in Wet H 2 S Environments” Definition of environment to be more susceptible to HIC SOHIC or blistering - process temp.: Ambient to 150 °C -H 2 S: 2000ppm + ph 7.8 -H 2 S: 50 ppm + ph 5 - presence of HCN + others Recommendations for repair: - Hardness of production welds 200 HB - Welding procedure qualification hardness 248 HV10 for HAZ and WELD - PWHT to be considered NACE RP0472 – 2000 2

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5 1 Requirements for steel plates in sour service Verfasser/Dokument 51 NACE MR0103 – 2003 „Materials Resistant to Sulfide Stress Cracking in Corrosive Petroleum Refining environments“ NACE MR0175: for oil- and gas handling systems NACE MR0103: for refinery service it based on the experience with MR0175 and other NACE publications. Specific Process Conditions: • 50 ppm H 2 S dissolved in H 2 O or if • pH 4 + some H 2 S or if • pH 7.6 + 20 ppm HCN + some H 2 S or if • 0.05 PSIA H 2 S in gas phase • Also reference to NACE RP0472 requirements

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5 2 Requirements for steel plates in sour service Verfasser/Dokument 52 NACE MR0103 – 2003 2 „Materials Resistant to Sulfide Stress Cracking in Corrosive Petroleum Refining environments“ Responsibility of the user: -HAZ -hardness - Residual stresses -R m increase risk increase Hardness Base Metal 22 HRC or also 248 HV 10 Cold deformation 5 otherwise stress relieved

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Revision 3 PV Plates for sour Service Verfasser/Dokument 53 Production of HIC-resistant steels

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5 4 Requirements for steel plates in sour service Verfasser/Dokument 54 How to produce HIC and SSC-resistant steel plates Basis: Well developed know how Dillinger Hütte GTS has been engaged in this field for more than 20 years Adequate production installations Permanent exchange with the endusers Follow up in international research projects

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5 5 Requirements for steel plates in sour service Verfasser/Dokument 55 DICREST-route hot metal desulphurisation deep vacuum degassing special chemical composition C Mn S P cleanliness stirring by Argon special casting parameter no bulging adapted superheating intensified QA-process special care to avoid unacceptable segregations high shape factor rolling strong reduction in thickness per rolling pass Requirements for homogeneous Dillinger Crack Resistant Steel plates

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5 6 Requirements for steel plates in sour service Verfasser/Dokument 56 Production route in the steel plant hot metal desulphur- isation dephosphorisation decarburisation denitrogenisation slag conditioning steel desulphur- isation removal of: Carbon Sulphur Nitrogen Hydrogen temperature adjustment Hot metal desulphurisation BOF converter Argon stirring process Degassing process Heating Casting cleanliness avoiding: - reoxidation - resulphurisation CaC Mg 2 O 2 Ar/N 2 Ar Ar Ar Ar O 2 analysis adjustment objective:

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5 7 Requirements for steel plates in sour service Verfasser/Dokument 57

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5 8 Requirements for steel plates in sour service Verfasser/Dokument 58 Curved caster Vertcal caster Distance from the fixed side in Total oxygen in ppm 10 20 30 40 50 60 70 80 90 20 40 60 80 100 0 0 Curved caster r 5.0 m v 1.0 m/min Vertical caster c v 0.5 m/min c Vertical caster Vertical caster Curved caster Inclusion distribution for different caster types

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5 9 Requirements for steel plates in sour service Verfasser/Dokument 59 Results from HIC test according to NACE TM 0284-96 for one single heat in dependence of the cast length of DICREST 15 pressure vessel steel test solution acc. to TM 0284-96: A pH3. 20 40 60 80 100 96 HIC resistant steel 75 of cast length cast length non sour gas non sour gas Frequency for CLR av. of 9 sections 15 end 0 begin Aspects of quality assurance: HIC properties and cast length

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6 0 Requirements for steel plates in sour service Verfasser/Dokument 60 Influence of High Shape Factor Rolling

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6 1 Requirements for steel plates in sour service Verfasser/Dokument 61 Optimized production steps for DICREST plates in the heavy plate mill

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6 2 Requirements for steel plates in sour service Verfasser/Dokument 62 Example of deviation in casting parameter combination incident risk range prohibited from release additional testing additional testing cast strand length position cracking extend in HIC test acceptance criteria Aspects of quality assurance: casting incidents

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6 3 Requirements for steel plates in sour service Verfasser/Dokument 63 Equipment: 8 laboratory fume hoods 7 for tests 1 for cleaning overall 39 connections for tests vessels 12 connections for SSC tensile tests CorTest rings equipped with computer aided monitoring of specimen failure 3 independent gas supply systems for parallel use of 3 different types of test gases temperature adjustment and control system Additionally health and safety-installations: gas detection systems flame guard system to maintain H 2 S combustion activated carbon filters in the exhaust air conduit collecting tanks for all waste waters from the process Test laboratory of Dillinger Hütte GTS to measure sour gas susceptibility:

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6 4 Requirements for steel plates in sour service Verfasser/Dokument 64 What can DH offer Sour service

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6 5 Requirements for steel plates in sour service Verfasser/Dokument 65 40mm 40 - 80mm 80mm Requested thicknesses Actual statistics of requested standards and thicknesses About 5 of the overall DICREST tonnage is requested in grades other than SA 516

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6 6 Requirements for steel plates in sour service Verfasser/Dokument 66 Note: Acceptance criteria are defined as the average of all sections of all specimens per plate ETC Extent of transverse cracking b max ELC Extent of longitudinal cracking a max 1 The requested test solution must be stated in the order in case of DICREST 15 Dillinger Hütte´s standardised offer for HIC resistant plates: DICREST CLR CTR CSR DICREST 5 80 mm A pH 3 ≤ 5 ≤ 1.5 ≤ 0.5 DICREST 10 80 mm A pH 3 ≤ 10 ≤ 3 ≤ 1 A pH 3 ≤ 15 ≤ 5 ≤ 2 B pH 5 ≤ 0.5 ≤ 0.1 ≤ 0.05 test solution acc. TM 0284-96 acceptance criteria DICREST 15 1 150 mm grade max. plate thickness 100 ⋅ ∑ W a CLR 100 ⋅ ∑ T b CTR 100 ⋅ ⋅ ⋅ ∑ T W b a CSR

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6 7 Requirements for steel plates in sour service Verfasser/Dokument 67 No. of Sections acc. to NACE TM0284-03 for t 88mm 15 No. of sections plate thickness CLR CTR CSR t ≤ 30mm 15 5 0.5 30mm t ≤ 40mm 15 3 0.5 40mm t ≤ 110mm 15 3 0.1 t ≤ 30mm 10 3 0.5 30mm t ≤ 110mm 10 2 0.1 t ≤ 15mm 5 1.5 0.5 15mm t ≤ 30mm 5 1.5 0.5 30mm t ≤ 110mm 5 1 0.1 1 3 9 resp. 15 Remark: All other requirements on request Actual acceptance levels for DICREST 5 plates in pH3 solution HIC tested in dependence on averaging the values for a certain no. of section

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6 8 Requirements for steel plates in sour service Verfasser/Dokument 68 Optimisation of CLR-values in NACE TM 0284-96 solution A through application of special DICREST-production route steel grades: A 516 Gr. 60 65 and 70 plate thickness 6-80 mm compared to Pseudo HIC-plates with a package of certain Pseudo-HIC measures. 0 10 20 30 40 50 60 70 80 90 2 2 ≤ 4 4 ≤ 6 6 ≤ 8 8 ≤ 10 10 ≤ 20 20 ≤ 40 40 HIC-resistant Pseudo-HIC Risk assessment on real HIC and Pseudo-HIC plates Percentage 50

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6 9 Requirements for steel plates in sour service Verfasser/Dokument 69

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7 0 Requirements for steel plates in sour service Verfasser/Dokument 70 DICREST ex mill and ex stock www.ancoferwaldram.nl

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7 1 Requirements for steel plates in sour service Verfasser/Dokument 71 thickness: 8 -80 mm grades SA 516 grade 60 65 or 70 delivery condition: normalised toughness requirements acc. SA20-S5 HIC testing frequency: per heat on the thinnest and thickest plate HIC test per NACE TM0284-2003 solution A pH3 hot tensile test at 400°C ultrasonic testing: acc. A578 ed. 2001 S 2.2 additionally: - conformity in harness and Ni-content to NACE MR0175 - banding check acc. to E 1268 once per heat for information DiME specification is more customized especially for Middle Eastern market in thickness range from 10 to 50 mm Specification details of DICREST stock plates AWS

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7 2 Requirements for steel plates in sour service Verfasser/Dokument 72 sour service becomes more and more important research and standardising efforts further ongoing 2 3 major failure mechanisms are important HIC SSC and probably SOHIC SSC rules have low influence on steel making practice. The phenomenon is mostly seen at hard HAZ or hard base metal. DH-GTS applies DICREST route. SOHIC is not quite fully understood. Most appearances are related to failures in HAZ no proper test method research is going on. Q+T steels show advantages. HIC resistant steels need a special manufacturing route and require a lot of experience know how Conclusion

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7 3 Requirements for steel plates in sour service Verfasser/Dokument 73 We contribute with 450.000 t of HIC resistant 1 linepipe and pv-plates per year 1 with certified HIC-resistance Dillinger Hütte GTS is prepared for the needs of sour service

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7 4 Requirements for steel plates in sour service Verfasser/Dokument 74 We can not transform sour to sweet... ... but we can help you take it with a smile

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2011 NACE CORROSION 2011 STG 32 – Advances in Materials for Oil Gas Production Mitigation of Sulfide Stress Cracking in Down Hole P110 Components via Low Plasticity Burnishing Jeremy Scheel Doug Hornbach Paul Prevéy Lambda Technologies 5521 Fair Lane Cincinnati OH 45227-3401 USA Darrel Chelette Peter Moore U.S. Steel Tubular Products Inc. 10343 North Sam Houston Park Drive 120 Houston TX 77064 USA ABSTRACT Sulfide stress cracking SSC along with hydrogen embrittlement HE prevents the use of less expensive high strength carbon steel alloys in the recovery of fossil fuels in H 2 S containing ‘sour’ service environments that are commonly seen in deep well fossil fuel recovery efforts. High magnitude tensile stresses are generated by connection interferences created during power make-up of down hole tubular components. When subject to service loads the stresses are increased further providing the high tensile stresses necessary for SSC initiation. Because these alloys processed into high strength grades are not suited for fully saturated sour service environments the current solution is to use or develop much more expensive alloys with increased corrosion-cracking resistance or limit their use to significantly weaker sour environments or higher operating temperatures. ©2011 by NACE International. Requests for permission to publish this manuscript in any form in part or in whole must be in writing to NACE International Publications Division 1440 South Creek Drive Houston Texas 77084. The material presented and the views expressed in this paper are solely those of the authors and are not necessarily endorsed by the Association. 1 Paper No. 11115

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Introduction of stable high magnitude compressive residual stresses into less expensive carbon steel alloys alleviates the tensile stresses and mitigates SSC while also improving fatigue strength. This could allow the potential of using less expensive alloys in sour environments. Low plasticity burnishing LPB is highly effective when applied to metallic components using a proven reproducible process of producing deep high magnitude compressive residual stresses in complex geometric components without altering the geometry design or chemistry. The LPB process applied with advanced control systems is presently being employed to treat components resulting in a substantial increase in service life through SSC mitigation and improved fatigue life. The benefits of LPB have been evaluated on full size specimens of uni-axial hoop stress loaded coupling blanks and C-ring specimens manufactured from quench and tempered API P110 grade steel with a yield strength of 132 ksi 910 MPa. Specimens were exposed to 100 NACE TM0177-Solution A at 1 bara H 2 S in both the LPB treated and untreated condition. The time to failure was documented along with the increase in life resulting from LPB treatment. LPB was successful in completely mitigating SSC in each test specimen up to 85 SMYS hoop tension and in each case met or exceeded the 720-hour exposure time defined in NACE TM0177. At an applied fiber stress of 90 SMYS the C-ring samples have exceeded exposures of 840 hours without failure. The initial results indicate that LPB processing of down hole tubular components may provide an alternative economical means of SSC mitigation and greatly reducing risk of component failure in sour environments. Key Words: low plasticity burnishing sulfide stress cracking fatigue residual stress sour service INTRODUCTION Surface enhancement of metals inducing a layer of surface compressive residual stresses in metallic components has long been recognized 1-4 to enhance fatigue strength and mitigate stress cracking. The fatigue strength of many engineering components is often improved by methods including rolling or shot peening. Modern surface enhancement treatments such as low plasticity burnishing LPB 5 laser shock peening LSP 6 and ultrasonic peening 7 have emerged that in varying degrees benefit fatigue and stress corrosion prone components. Maximum benefits are obtained when deep compression is achieved with minimal cold working of the surface. Environmentally assisted cracking EAC in the form of Sulfide Stress Cracking SSC Stress Corrosion Cracking SCC and Hydrogen Embrittlement HE prevent the use of less expensive high strength carbon steel alloys in the recovery of fossil fuels in corrosive-cracking environments commonly seen in offshore and deep well recovery efforts. Tensile residual stresses generated from straightening machining and connection make-up when added to applied stresses during down hole operations in high-pressure environments are significant contributors to EAC and fatigue failure. Because these alloys at high strength levels are not suited for sulfide or chloride environments the current solution is to use or create much more expensive alloys with increased corrosion-cracking resistance to mitigate the problems or limit their use to significantly weaker sour environments. Introducing compressive residual stresses into less expensive carbon steel alloys can dramatically reduce the risk of failure mitigate SSC HE and SCC and improve fatigue 2

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strength. 8 This could allow the potential of using less expensive alloys in harsh environments where they currently are unable to be used. Low plasticity burnishing LPB is highly effective reliable and reproducible method of producing deep compressive residual stresses in complex geometric components. With advanced control systems LPB can be applied using a closed loop feedback surface enhancement method capable of introducing a customized compressive residual stress field specifically tailored for each application. LPB so applied produces a very smooth surface finish which aids in nondestructive inspection and examination. LPB tooling can be integrated with existing equipment used for manufacture and repair of down hole tubular products. SSC susceptibility in high strength API P110 grade tubular products prevents their use in 100 H 2 S sour environments at temperatures less than 79°C 175 ° F. 910 As more deep wells and offshore resources are probed and recovered it is imperative to mitigate the problem of EAC in a cost effective manner. Laboratory testing in standard 100 NACE TM0177-Solution A 11 at 1 bara H 2 S liquid environment has shown that the LPB process can be employed to treat components with the effect of providing a substantial increase in service life and SSC mitigation. The LPB technology is now being evaluated to determine if it can play a pivotal role in creating more reliable and efficient fossil fuel recovery systems that are capable of safely and reliably operating in aggressive environments. LPB processing has successfully been used to mitigate EAC in high strength steel stainless steels and aluminums used in both the aerospace and nuclear industries. LPB technology was developed in conjunction with NASA’s SBIR program and is currently used in production of parts used in the aerospace medical and nuclear industries and on many different metal alloys. 12-16 EXPERIMENTAL PROCEDURE Material C- ring specimens and full size coupling blank specimens were sectioned from a length of API P110 grade quench and tempered coupling stock. Figure 1 shows the coupling stock used to manufacture the specimens. Figure 2 shows an example of each of the 2 geometries tested in this investigation. Figure 1: API P110 Quench + Temper coupling stock. 3

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Figure 2: Examples of tested geometries: A C-ring specimen and B full sized coupling blank in test fixture. Specimen Processing LPB process parameters were developed to achieve nominally 0.040 in. 1 mm depth of compression. Samples were processed on a CNC mill or lathe to allow positioning of the LPB tool in a series of passes along the region to be processed while controlling the burnishing pressure to develop the pre-determined magnitude of compressive stress with controlled low cold working. The full lengths of the outer diameter of the full sized coupling blanks were LPB processed. The C-ring specimens were processed on the exposed section of the outside diameter. The LPB process has been previously documented in detail. 17 X-ray Diffraction Residual Stress Analysis X-ray diffraction residual stress measurements were made at the surface and at several depths below the surface on the outside diameter of both LPB and untreated specimens to characterize the residual stress distributions. Measurements were made in the axial direction employing a sin 2 ψ technique and the diffraction of chromium K α1 radiation from the 211 planes of steel. Material was removed electrolytically for subsurface measurement in order to minimize possible alteration of the subsurface residual stress distribution. The measurements were corrected for both the penetration of the radiation into the subsurface stress gradient and for stress relaxation caused by layer removal. The value of the x-ray elastic constants required to calculate the macroscopic residual stress from the strain normal to the 211 planes of steel were determined in accordance with ASTM E1426-9. 1819 Systematic errors were monitored per ASTM specification E915. 20 Surface Roughness The improvement in surface roughness was documented for LPB vs. un-treated coupling material. Surface roughness measurements were performed on both the untreated and LPB treated coupling blanks using a standard surface roughness tester. The Ra surface roughness was calculated over a 0.50 in. 12.7 mm evaluation length in the axial direction. A B 4

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SSC Testing SSC Testing was conducted on 4 ½ in. 114.3 mm API P110 quench and temper coupling stock 5 in. 127 mm outside diameter. The coupling stock was sectioned into C-Ring specimens per NACE TM0177-Method C for testing and the full sized coupling stock blanks were machined to create an inside diameter of 4.375 in 111.1 mm and provide a sealing surface for the internal pressure seals. Specimens were tested in both the un-treated condition as well as after LPB processing to determine the differential effects resulting from LPB treatment. C-ring Testing: Testing was performed on LPB treated and un-treated C-ring specimens per NACE TM0177- Method C. All testing was performed in 100 NACE TM0177-Solution A at 1 bara H 2 S at 25° C the pH was monitored continuously throughout testing to ensure conformance to NACE TM0177. Specimens were sectioned from a length of API P110 coupling stock. The specimens were loaded initially to nominally 45 of SMYS. After exposure to at least 720 hours the specimens were tested at 80 85 and 90 of SMYS. Stress on the specimens was monitored continuously using strain gage rosettes placed on the inner diameter opposite the exposed location of maximum applied tension. The entire specimen except the outer gage region was coated in a polymer based stop off coating after loading and prior to immersion in solution. Figure 3 shows a C-ring specimen ready for testing. Figure 3: C-ring specimen prior to testing. Full Sized Pressurized Coupling blank Testing: Testing of full size coupling blanks was performed using a custom made holding fixture connected to a pressurizing test station. Specimens were tested in both the LPB treated and un-treated conditions. All tests were performed in 100 NACE TM0177-Solution A at 1 bara H 2 S at room temperature with the pH continuously monitored. The full sized coupling blanks were internally pressurized hydraulically to impart the desired amount of applied hoop stress. Test solution was monitored for pH and refreshed as needed to conform to the NACE TM0177 standard. Testing was conducted until specimen failure or a run out life of 720 hours 30 days or more was achieved per NACE TM0177 standard. Specimens were tested at 45 80 and 5

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85 of SMYS. Pressure was monitored continuously throughout the test and a timer was placed in the circuit to trip upon sample failure. Figure 4: Full size coupling blank pressurized test apparatus and setup. RESULTS AND DISCUSSION X-ray Diffraction Residual Stress Analysis X-ray diffraction residual stress vs. depth results for untreated and LPB processed API P110 coupling blanks are presented graphically in Figure 5. Compressive stresses are shown as negative values and tensile stresses as positive in units of ksi 10 3 psi and MPa 10 6 N/m 2 . Compared to the untreated condition LPB produced a compressive residual stress field with a much greater magnitude of compression 10X and over 2X the depth of compression. The magnitude of compression is near the SMYS of 110 ksi 759 MPa at the surface. LPB produces much less cold working than conventional processes and ensures the deep fiber layers remain in stable compression even at high temperature or in the case of mechanical overload as has been demonstrated in prior work. 21 6

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0 1020 304050 -120 -100 -80 -60 -40 -20 0 20 Untreated Quench + Temper LPB Residual Stress ksi Depth x 10 -3 in. 0 200 400 600 800 1000 1200 -800 -700 -600 -500 -400 -300 -200 -100 0 100 LPB Quench + Temper Untreated Residual Stress MPa Depth x 10 -3 mm Figure 5: Residual stress comparison for LPB processed and un-treated material. Surface Roughness The improvement in surface roughness after LPB processing was quantified using the Ra surface roughness. LPB improved the surface finish by a factor of 2.6X. This can aid in NDI examination as well as reduce friction in service. Figure 6 displays the results graphically. Each value is an average of three repeat measurements. 7

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0 50 100 150 200 250 300 Surface Roughness - Ra μm Average of 3 Measurements 241.7 92.0 AVERAGED SURFACE ROUGHNESS API P110 Steel Coupling UNTREATED Surface Roughness - Ra μ in LPB PROCESSED 0 1 2 3 4 5 6 7 Figure 6: Surface roughness for LPB processed and un-treated API P110 coupling blank. SSC Testing The SSC testing data is presented graphically below in Figures 7 8. The un-treated C-ring specimen with an OD exposed surface failed in 10 hours at a stress of 45 SMYS The LPB processed specimens exceeded the run-out life of 720 hours at 45 80 85 and 90 of SMYS exceeding typical hold-time requirements for testing in a sour service environment. The full sized coupling blank test results are very similar with the un-treated coupling blank failing after the entire OD surface was exposed for 37.5 hours. The LPB processed specimens exceeded 720 hours at 45 80 and 85 SMYS stress levels while surpassing typical hold- time requirements for sour service testing. The second full sized LPB coupling blank ran for a total of 1454.75 hours in solution before testing was terminated and the specimen was removed from solution for dye inspection which revealed no cracking. These test results demonstrate the dramatic improvement achieved by the LPB treatment compared to the un- treated P110 material. A macro photo comparison of a failed untreated C-ring specimen and a run out LPB C-ring specimen is shown in Figure 9. Dye penetrent was used to reveal the axial SSC failure in the un-treated coupling blank shown in Figure 10. Figure 11 shows the LPB coupling blank after timed run out at 85 SMYS with and without FDI developer documenting that there are no cracks of any size beginning to initiate on the specimen. The SSC testing results show definitively that LPB is able to mitigate SSC cracking in common API P110 steel and dramatically increase the life. The 85 SMYS stress level is regarded as an aggressive performance test for metal that is in direct contact with a 100 H 2 S saturated environment. API Specification 5CT uses 80 SMYS stress levels for C90 and T95 tensile specimens and the next edition will likely add 85 SMYS stress level for a new C110 grade. 8

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0 200 400 600 800 1000 1200 1400 1600 1800 LPB PROCESSED 90 SMYS RUN-OUT EXCEEDED NACE TM0177 841 Hours RUN-OUT EXCEEDED NACE TM0177 820 Hours LPB PROCESSED 85 SMYS TIME Hours RUN-OUT EXCEEDED NACE TM0177 822 Hours NACE TM0177 RUN-OUT 720 Hrs FAILED 10 Hrs UNTREATED Quench +Temper 45 SMYS LPB PROCESSED 45 SMYS LPB PROCESSED 80 SMYS RUN-OUT EXCEEDED NACE TM0177 by 2X 1719 Hours Test Stopped API P110 STEEL C-RING TESTING NACE A Solution 25° C Figure 7: C-ring testing results. 0 100 200 300 400 500 600 700 800 1454.75 hrs Total Time Exposed LPB PROCESSED 85 SMYS RUN - OUT 720.25 hrs TIME Hours RUN - OUT 734.5 hrs NACE TM0177 RUN-OUT 720 Hrs FAILED 37.5 Hrs UNTREATED Quench+Temper 45 SMYS LPB PROCESSED 80 SMYS API P110 STEEL COUPLING PRESSURE TEST NACE A Solution 25° C Figure 8: Full Sized Coupling blank Test Results. 9

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Figure 9: Comparison of LPB treated and un-treated C-ring specimens after testing. The un- treated specimen failed in 10 hours at 45 of SMYS. The LPB specimens ran-out at 45 80 85 and 90 SMYS with no cracking. Figure 10: FDI inspection of failed un-treated coupling blank revealing thru wall axial SSC. Figure 11: LPB processed coupling blank and test fixture after run-out at 85 SMYS. FDI developer showing no signs of crack initiation. AXIAL FAILURE AXIAL FAILURE LPB UNTREATED FAILED 10

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CONCLUSIONS • LPB imparted a deep compressive layer of stable residual compression over 2X deeper and 10X greater in magnitude than the untreated coupling blanks. • LPB was able to completely mitigate SSC failure in all tested specimens. The full sized coupling blank test exceeded the NACE TM0177 720 hour NACE A exposure time requirement at 45 80 and 85 of the SMYS of 110 ksi 759 MPa. • The LPB processed C-ring tests exceeded NACE TM0177 time requirements at stresses levels equal to 45 80 85 and 90 of SMYS. • The untreated coupling blanks and c-ring specimens failed in 33 hrs and 10 hrs respectively at a stress load of 45 SMYS. • Use of an engineered deep compressive stress field using LPB to mitigate SSC was successful on API P110 quench and temper coupling blank specimens. REFERENCES 1. Frost N.E. Marsh K.J. Pook L.P. 1974 Metal Fatigue Oxford University Press. 2. Fuchs H.O. and Stephens R.I. 1980 Metal Fatigue In Engineering John Wiley Sons. 3. Berns H. and Weber L. 1984 "Influence of Residual Stresses on Crack Growth" Impact Surface Treatment edited by S.A. Meguid Elsevier 33-44. 4. Ferreira J.A.M. Boorrego L.F.P. and Costa J.D.M. 1996 "Effects of Surface Treatments on the Fatigue of Notched Bend Specimens" Fatigue Fract. Engng. Mater. Struct. Vol. 19 No.1 pp 111-117. 5. Prevéy P.S. Telesman J. Gabb T. and Kantzos P. 2000 “FOD Resistance and Fatigue Crack Arrest in Low Plasticity Burnished IN718” Proc of the 5 th National High Cycle Fatigue Conference Chandler AZ. March 7-9. 6. Clauer A.H. 1996 "Laser Shock Peening for Fatigue Resistance" Surface Performance of Titanium J.K. Gregory et al Editors TMS Warrendale PA pp 217-230. 7. T. Watanabe K. Hattori et al 2002 “Effect of Ultrasonic Shot Peening on Fatigue Strength of High Strength Steel” Proc. ICSP8 Garmisch-Partenkirchen Germany Ed. L. Wagner pg 305- 310. 8. Paul S. Prevéy N Jayaraman "Overview of Low Plasticity Burnishing for Mitigation of Fatigue Damage Mechanisms" Proceedings of ICSP 9 Paris Marne la Vallee France Sept. 6-92005. 9. Snape E.: “Sulfide Stress Corrosion of Some Medium and Low Alloy Steels” Corrosion June 1967 23 326-332. 10. Carter C.S. and Hyatt M.V.: “Review of Stress Corrosion Cracking in Low Alloy Steels With Yield Strength Below 150 ksi” SCC and Hydrogen Embrittlement of Iron Base Alloy NACE Reference Book No. 5 1977 524-600. 11. NACE Standard TM0177-2005: Laboratory Testing of Metals to Specific Forms of Environmental Cracking NACE International. 12. J. Scheel D. Hornbach P. Prevey “Mitigation of Stress Corrosion Cracking in Nuclear Weldments Using Low Plasticity Burnishing” Proceedings of the 16 th International Conference on Nuclear Engineering ICONE16 May 11-15 2008 Orlando FL. 13. N. Jayaraman P. Prevéy “An Overview of the use of Engineered Compressive Residual Stresses to Mitigate SCC and Corrosion Fatigue” Proceedings of 2005 Tri-Service Corrosion Conference Orlando FL Nov. 14-18 2005. 14. D.H. Hornbach and P.S. Prevéy “Tensile Residual Stress Fields Produced in Austenitic Alloy Weldments” Proceedings: Energy Week Conference Book IV Jan. 28-30 Houston TX ASME International 1997. 11

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15. P.S. Prevey et al. “Effect of Prior Machining Deformation on the Development of Tensile residual Stresses in Weld Fabricated Nuclear Components” Journal of Materials Engineering and Performance vol. 51 Materials Park OH ASM International 1996 pp. 51-56. 16. D. Hornbach P. Prevéy “Reducing Corrosion Fatigue and SCC Failures in 300M Steel Landing Gear Using Low Plasticity Burnishing” Proceedings of 2007 SAE AeroTech Congress Exhibition Los Angeles CA September 17-20 2007. 17. P. Prevey. “Burnishing Method and Apparatus for Providing a Layer of Compressive Residual Stress in the Surface of a Workpiece.” US Patent 5826453 Oct. 27 1998. 18. Cullity B.D. 1978 Elements of X-ray Diffraction 2nd ed. Reading MA: Addison-Wesley pp. 447-476. 19. Prevéy P.S. 1986 “X-Ray Diffraction Residual Stress Techniques” Metals Handbook 10 Metals Park OH: ASM pp 380-392. 20. ASTM Standard E915 2010 "Standard Test Method for Verifying the Alignment of X-Ray Diffraction Instrumentation for Residual Stress Measurement" ASTM International West Conshohocken PA 2003 DOI: 10.1520/E0915-10 www.astm.org. 21. Paul S. Prevéy “The Effect of Cold Work on the Thermal Stability of Residual Compression in Surface Enhanced IN718” Proceedings of the 20 th ASM Materials Solutions Conference and Exposition St. Louis MO Oct. 10-12 2000. 12

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Hydrogen Induced Damage in Pipeline Steels by Garrett R. Angus

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Copyright by Garrett R. Angus 2014 All Rights Reserved

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ii A thesis submitted to the Faculty and the Board of Trustees of the Colorado School of Mines in partial fulfillment of the requirements for the degree of Masters of Science Metallurgical and Materials Engineering. Golden Colorado Date: ______________ Signed: _________________________ Garrett R. Angus Signed: _________________________ Dr. Kip O Findley Thesis Advisor Golden Colorado Date: ______________ Signed: _________________________ Dr. Chester J. Van Tyne FIERF Professor and Interim Department Head Department of Metallurgical and Materials Engineering

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iii Abstract The hydrogen induced cracking HIC resistance of several grades of plate steels was investigated using electrolytic hydrogen charging. HIC generated by electrolytic charging was also compared to the industrial standard test for HIC the NACE standard TM0284. The electrolytic charging EC apparatus was designed to optimize the reproducibility of the HIC results and the robustness of the components during long charging times. A characterization study on the EC apparatus was undertaken. Alterations to applied current density and charging time were conducted on a highly susceptible plate steel 100XF to assess HIC damage as a function of charging conditions. Intermediate current densities of 10 to 15 mA/cm 2 produced the greatest extent of cracking without significant corrosion related surface damage. The hydrogen charging time did not greatly affect the extent and depth of cracking for test times between 24 to 48 hours. Thus for subsequent experiments the applied current density was set to 15 mA/cm 2 and the charging time was set to 24 hours. Plate steel grades X52 X60 X70 and 100XF were prestrained in tension to various levels and then electrolytically charged with hydrogen or tested with the NACE standard TM0284 test solution A saturated with H 2 S g to induce HIC. Prestrain was introduced to assess its impact on HIC. Hydrogen damage was quantified with the crack ratios defined in the NACE Standard TM0284. The results from the EC and NACE methods were very comparable to one with respect to the magnitude of cracking and the trends between alloy and pre-strain conditions observed. Both methods showed that HIC substantially increased for the high strength 100XF steel compared to the lower strength alloys. This is consistent with NACE recommendations for HIC resistance steels which suggests that alloy strength should be less than 116 ksi 800 MPa or 248 HV 22 HRC. The HIC results were largely independent of the pre-strain levels imposed within the statistical accuracy of the evaluation method employed. The total irreversibly trapped and diffusible hydrogen amounts were measured or estimated for each condition using a LECO interstitial analyzer and the American Welding Society method for measuring diffusible hydrogen concentrations. The total amount of diffusible hydrogen was highest for the 100XFalloy and lowest for the X52 alloy. The amount of trapped hydrogen was similar for all the alloys implying that the number of irreversible trap sites were comparable. However the diffusible hydrogen content was greatest for the 100XF alloy and lowest for the X52 alloy which is believed to be related to the relatively high amount of grain boundary area and high dislocation density of the 100XF alloy. A qualitative analysis on the effect of microstructure and nonmetallic inclusions on HIC was performed and produced results that confirmed findings from literature. Cracking was observed around nonmetallic inclusions such as sulfides and oxides in the metal matrix. For materials in which both inclusion types were present X60 and X70 HIC originated and was observed most often around sulfide type inclusions.

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iv Table of Contents Abstract ........................................................................................................................................................ iii Table of Contents ......................................................................................................................................... iv LIST OF FIGURES .......................................................................................................................................... vi LIST OF TABLES .......................................................................................................................................... xiii Acknowledgements .................................................................................................................................... xiv CHAPTER 1: Introduction ....................................................................................................................... 1 1.1 Research Objectives ...................................................................................................................... 1 CHAPTER 2: Background ....................................................................................................................... 3 2.1 Hydrogen Entry into Steel and Hydrogen-Assisted Cracking Mechanisms ................................. 3 2.2 Metallurgical Variables ................................................................................................................. 5 2.2.1 Deoxidation and Casting Techniques.................................................................................... 6 2.2.2 Microstructural Variables ..................................................................................................... 8 2.2.3 Mechanical Processing Effects ........................................................................................... 11 2.3 Experimental Methods to Test Plate Steels for HIC Susceptibility ............................................ 13 2.3.1 NACE Standard Test TM0284 ............................................................................................ 13 2.3.2 Electrolytic Charging .......................................................................................................... 15 CHAPTER 3: Experimental Design and Methods ................................................................................. 17 3.1 Experimental Design ................................................................................................................... 17 3.2 Experimental Materials and Methods ......................................................................................... 17 3.2.1 Materials ............................................................................................................................. 17 3.2.2 Mechanical Properties ......................................................................................................... 18 3.2.3 Microstructural Analysis ..................................................................................................... 18 3.2.4 Introduction of Prestrain ..................................................................................................... 19 3.2.5 Hardness Traverse and Determination of Microstructural Dependence on Plate Thickness ……………………………………………………………………………………………..21 3.3 HIC Sample Generation and Preparation for HIC Testing ......................................................... 21 3.4 NACE Standard TM0284 H 2 S Method .................................................................................... 22 3.5 Electrolytic Charging Methodology EC ................................................................................... 26 3.5.1 Initial Design and Fabrication of New Charging Apparatus ............................................... 27 3.5.2 Characterization Study and Experimental Procedure for EC .............................................. 30

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v 3.6 Assessment of Hydrogen Damage .............................................................................................. 32 3.7 LECO® Hydrogen Analysis ....................................................................................................... 35 3.7.1 LECO® Hydrogen Analysis Sample Generation................................................................ 35 3.7.2 LECO® Hydrogen Analysis Experimental Procedure ........................................................ 37 3.8 Diffusible Hydrogen Content Determined by Mercury Displacement ....................................... 37 3.8.1 Mercury Displacement Sample Generation ........................................................................ 37 3.8.2 Mercury Displacement Experimental Procedure ................................................................ 38 CHAPTER 4: Results and Discussion.................................................................................................... 43 4.1 Microstructure and Nonmetallic Inclusions ................................................................................ 43 4.2 Mechanical Properties ................................................................................................................. 49 4.3 Characterization Study ................................................................................................................ 54 4.4 HIC Susceptibility Measurements .............................................................................................. 58 4.4.1 Influence of Prestrain on Calculated Critical Crack Ratios ................................................ 59 4.4.2 Influence of Mechanical Properties on HIC ........................................................................ 62 4.5 Hydrogen Analysis ...................................................................................................................... 63 4.5.1 Average Hydrogen Contents Total and Trapped from LECO® Analysis .......................... 64 4.5.2 Average Diffusible Hydrogen Content from LECO® and Mercury Displacement Analysis ……………………………………………………………………………………………..66 4.5.3 Influence of Sample Location on Diffusible Hydrogen Results Sample Size Effect ....... 67 4.5.4 HIC Susceptibility and Diffusible and Trapped Hydrogen Contents from LECO® Analysis ……………………………………………………………………………………………..70 4.6 Evaluation of HIC with respect to Microstructure and Nonmetallic Inclusions ......................... 71 CHAPTER 5: Summary and Conclusions .............................................................................................. 79 5.1 Electrolytic Charging .................................................................................................................. 79 5.2 HIC Susceptibility Measurements .............................................................................................. 79 5.2.1 Influence of Mechanical Properties .................................................................................... 79 5.2.2 Influence of Uniaxial Prestrain ........................................................................................... 80 5.3 Hydrogen Analysis ...................................................................................................................... 80 5.4 Evaluation of HIC with respect to Microstructure and Nonmetallic Inclusions ......................... 80 CHAPTER 6: Future Work .................................................................................................................... 81 References Cited ......................................................................................................................................... 82

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vi LIST OF FIGURES Figure 2.1 Schematic representation of the dissociation of hydrogen sulfide gas at the metal/solution interface and subsequent diffusion into the metal to areas where hydrogen-assisted cracking can occur 5. ...... 4 Figure 2.2 Schematic representation of enhanced plastic flow mechanism for hydrogen embrittlement. Parts A B and C show various configurations of inclusion size/elongation distribution and amount and the resulting morphology of cracks 3. ........................................................................................................ 5 Figure 2.3 Schematic representation of the influence of ingot material location on HIC susceptibility of large sized ingots used to produce hot-rolled product 3. ............................................................................... 7 Figure 2.4 Schematic diagram illustrating the hydrogen penetration through two varying degrees of banding in a ferrite/pearlite microstructure: a lower degree of banding and b higher degree of banding. In a the hydrogen penetration is hindered by pearlite along the ferrite regions. Adapted from 19. .................. 9 Figure 2.5 Examples of nonmetallic inclusions that form in fully killed steels. Shown are the changes of morphology due to rolling operations after casting. Adapted from 25. .............................................. 10 Figure 2.6 HIC susceptibility crack length ratio ratio as a function of cold rolling reduction. Lattice strain was determined through X-ray diffraction after cold rolling operations denoted as line broadening in the above figure. HIC testing was performed with respect to the NACE standard test TM0284. Steel used for HIC testing was a 0.07-C 1.22-Mn 0.006-S all values in wt pct micro-alloyed with Nb and V heavily controlled rolled steel. Adapted from 20. .............................................................................. 12 Figure 2.7 Effects of cold reduction at levels greater than 10 pct. Steels A and B have similar carbon equivalents 0.31 and were produced from an ingot whereas steel C was produced by continuous casting and has a higher carbon equivalent of 0.39. Steels A and C would be categorized are HIC-Resistant steels and steel B is a low sulfur grade steel. See text for more detailed chemical compositions of Steels A B and C. Adapted from 20. .................................................................................................................... 13 Figure 2.8 Schematic illustration of the two reactions for hydrogen that occur at the metal/solution interface. The reaction a occurs at a much higher rate than b. Reaction kinetics of a are suppressed in sour service applications due to the presence of sulfur ions Equation 2.3. Reaction b causes subsequent HIC within the material through hydrogen interactions with features in the metal matrix 36. .......... 14 Figure 3.1 Schematic of strips that were machined in the transverse direction with respect to the rolling direction and then subsequently prestrained to target levels. The prestrain and rolling directions are labeled. Strip length was dependent on the plate material that was machined. The length ranged from 450 mm for the X52 to 700 mm for the 100XF. ................................................................................................. 20 Figure 3.2 Example of a strip being prestrained using uniaxial tension. The entire length between grips was considered the gauge length. This gauge length was used in correspondence with target prestrain levels to set a predetermined crosshead displacement to achieve the desired prestain amount. ........... 20

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vii Figure 3.3 Test specimen geometry and orientation used in the NACE Standard and electrolytic charging studies all dimensions are in millimeters. For specimens that were subjected to electrolytic charging a taper hole was drilled and reamed in the transverse top face of the test specimen to allow for electrical conductivity. ......................................................................................................................................... 22 Figure 3.4 a NACE Standard TM0284 test vessel used to evaluate materials suscpetibilty to HIC b carousel used to hold test specimens allowing for adequate space for gas purge during testing. ....................... 26 Figure 3.5 Prior generation cell used for electrolytic charging by B. Rosner 7 and G. Angus 44. ................... 28 Figure 3.6 From left to right the images display the salient features for a the construction of the cathode electrode connection to be used in correspondence with b the steel specimens. The cathode electrode connection in a employs a taper pin fastened to a 316 SS rod. A corresponding taper hole is reamed into the top transverse face of the steel specimen. The cathode electrode connection is then inserted into the hole and frictionally locked with the test specimen. To protect this electrical connection the region around the connection is covered in epoxy. ............................................................................... 29 Figure 3.7 Schematic diagram for dual cell used for electrolytic charging experiments. ...................................... 29 Figure 3.8 Electrolytic hydrogen charging EC test apparatus with components labeled. Electrolytic charging EC was accomplished through electrochemical polarization using this test apparatus. Tests apparatus was designed to optimize the reproducibility of the HIC results and the stability of the components during long charging times ................................................................................................................... 31 Figure 3.9 Test specimen geometry and orientation used in the NACE Standard and electrolytic charging studies all dimensions are in millimeters a Faces 2 4 and 6 are used in the NACE Standard evaluation while Faces 1 - 6 are used for the electrolytic charging method. Face 1 is the furthest from the electrode link. b Schematic showing how cracks are measured on the evaluated faces 34. ............ 32 Figure 3.10 ImageJ interface that allows the user to import selected images into the software for analysis. Once imported the plate thickness is measured in pixels shown by the white line. Knowing this distance in pixels the “set scale” feature in the software allows the user to set the distance in pixels to a known distance. For this image and cross-section 502 pixels equaled 12.42 mm. The “set scale” feature also displays the resolution of the image relative to the scale. This resolution was used to determine the accuracy of the measurements. ............................................................................................................. 33 Figure 3.11 Example cracked face showing how images were used to a identify cracks that are offset from one another. b Example cracked face showing the measurement offset to determine if multiple cracks should be considered a single cracks separated by less than 0.5 mm were considered as a single crack 34. c The final measurement of the single cracks. Each crack in b was enhanced using the “zoom” feature to produce the images in c to aid in the precise measurement of the single crack. The white lines in c represent the length a and thickness b measurements used to calculate critical crack ratios. ........................................................................................................................................... 36

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viii Figure 3.12 “Daisy Chain” assembly components. From left to right the images display the salient features the five sub-sized specimens and copper wire of the component assembly leading to the assembly of the “Daisy Chain” to allow for simultaneous charging of multiple specimens for hydrogen content measurements. ...................................................................................................................................... 35 Figure 3.13 a Schematic representation with specimen geometry and the locations where test specimens were sectioned. The black areas represent material that was left after the sectioning process was complete. b Visual representation of samples fabricated for diffusible hydrogen analysis using EC. The sectioned areas allow for easy detachment after EC has been conducted. ............................................ 38 Figure 3.14 a Dimensions of the eudiometer used in the mercury displacement method to determine diffusible hydrogen amount b schematic representation of the use of a mercury-filled eudiometer to capture and measure the amount of diffusible hydrogen in the sample that is placed into the assembly. Adapted from 45. ............................................................................................................................... 40 Figure 3.15 Test setup for the mecury displacement method with the compontents labeled. .................................. 41 Figure 4.1 Secondary electron micrographs taken with the FESEM of X52 plate steel. Images from three different planes a transverse b longitudinal and c normal plane are shown. 2 pct Nital etch. ..... 44 Figure 4.2 Presence of degenerated pearlite in the X52 plate material. Image taken with the FESEM. 2 pct nital etch. ...................................................................................................................................................... 44 Figure 4.3 Secondary electron micrographs taken with the FESEM of X60 plate steel. Images from three different a transverse b longitudinal and c normal plane planes are shown. 2 pct nital etch. Black circles on each micrograph indicate the presence of a secondary microconstituent. .................. 45 Figure 4.4 Secondary electron micrographs taken with the FESEM of X70 plate steel. Images from three different planes a transverse b longitudinal and c normal plane are shown. 2 pct nital etch. Black circles on each micrograph indicate the presence of a secondary microconstituent. .................. 45 Figure 4.5 Secondary electron micrographs taken with the FESEM of 100XF plate steel. Images from three different planes a transverse b longitudinal and c normal plane are shown. 2 pct nital etch. Black circles on each micrograph indicate the presence of a secondary microconstituent. .................. 45 Figure 4.6 Nonmetallic inclusions observed in the four plate steels. Elements present in each image were confirmed by EDS mapping of the image shown. a Al-Ca-O X52 b Al-O X60 c Al-Mg-O-Ca-S X70 d Al-Mg-O 100XF see pdf version for color. ......................................................................... 46 Figure 4.7 Ternary diagrams in wt pct of two inclusions families a Al-Ca-S and b Ca-Mn-S present in the X52 plate steel. Ternary diagrams were generated by automatic feature analysis at the Nucor Indiana facility. Dashed regions 1and 2 in a and b respectively show the grouping of the relative compositional distribution of inclusions evaluated by AFA in reference to the three elements found in each ternary diagram. ............................................................................................................................ 47

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ix Figure 4.8 Ternary diagrams in wt pct of two inclusions families a Al-Ca-S and b Ca-Mn-S present in the X60 plate steel. Ternary diagrams were generated by automatic feature analysis at the Nucor Indiana facility. Dashed regions 1and 2 in a and b respectively show the grouping of the relative compositional distribution of inclusions evaluated by AFA in reference to the three elements found in each ternary diagram. Regions 3 in b indicate the presence of MnS type inclusions identified through AFA. ........................................................................................................................................ 48 Figure 4.9 Ternary diagrams in wt pct of two inclusions families a Al-Ca-S and b Ca-Mn-S present in the X70 plate steel. Ternary diagrams were generated by automatic feature analysis at the Nucor Indiana facility. Dashed regions 1and 2 in a and b respectively show the grouping of the relative compositional distribution of inclusions evaluated by AFA in reference to the three elements found in each ternary diagram. Regions 3 in b show the presence of MnS type inclusions identified through AFA. ..................................................................................................................................................... 48 Figure 4.10 Ternary diagrams in wt pct of two inclusions families a Al-Ca-S and b Ca-Mn-S present in the 100XF plate steel. Ternary diagrams were generated by automatic feature analysis at the Nucor Indiana facility. Dashed regions 1and 2 in a and b respectively show the grouping of the relative compositional distribution of inclusions evaluated by AFA in reference to the three elements found in each ternary diagram. Regions 3 in a show a region where the types of inclusions within the area if evaluated with other elements i.e. Fe and O the composition of the inclusion would be found to be different. Speculated that these inclusions would be Fe-S and Al-O type inclusions. .......................... 49 Figure 4.11 a Microhardness data from each rolled face to the center of the X52 plate thickness along with the average microhardness solid line with 90 pct confidence interval b etched 2 pct Nital and c non-etched macrographs taken from the edge to middle thickness. Cross-sections of the full plate thickness of the transverse plane with respect to rolling direction were evaluated. Microhardness measurements were taken every 0.254 mm 0.01 in 500 gmf load and 10 s dwell time. .................. 50 Figure 4.12 a Microhardness data from each edge to the center of the X60 plate thickness along with the average microhardness solid line with 90 pct confidence interval b etched 2 pct Nital and c non-etched macrographs taken from the edge to middle thickness. Cross-sections of the full plate thickness of the transverse plane with respect to rolling direction were evaluated. Microhardness measurements were taken every 0.254 mm 0.01 in 500 gmf load and 10 s dwell time. .................. 51 Figure 4.13 a Microhardness data from each edge to the center of the X70 plate thickness along with the average microhardness solid line with 90 pct confidence interval b etched 2 pct Nital and c non-etched macrographs taken from the edge to middle thickness. Cross-sections of the full plate thickness of the transverse plane with respect to rolling direction were evaluated. Microhardness measurements were taken every 0.254 mm 0.01 in 500 gmf load and 10 s dwell time. .................. 52

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x Figure 4.14 a Microhardness data from each edge to the center of the 100XF plate thickness along with the average microhardness solid line with 90 pct confidence interval b etched 2 pct Nital and c non-etched macrographs taken from the edge to middle thickness. Microhardness measurements were taken every 0.254 mm 0.01 in 500 gmf load and 10 s dwell time. ................................................... 53 Figure 4.15 Surface conditions of 100XF after cathodic charging at various applied current densities while keeping the duration of the test constant at 24 hours. a 0.80 b 1.50 c 5.0 d 10.0 e 15 and f 25 mA/cm 2 . Image taken using light optical flash photography. .......................................................... 55 Figure 4.16 Hydrogen Induced Cracks produced through electrolytic charging of 100XF at a current density of 15 mA/cm2 for 24 hours aFace 2 b Face 4. ..................................................................................... 57 Figure 4.17 Variation of calculated crack ratios versus applied current density for 100XF that was cathodically charged for 24 hours. CSR – Critical Size Ratio CLR – Critical Length Ratio and CTR – Critical Thickness Ratio. ................................................................................................................................... 57 Figure 4.18 Variation of average calculated crack ratios versus charging time for 100XF cathodically charged at an applied current density of 15 mA/cm 2 . ............................................................................................. 58 Figure 4.19 Critical Crack Ratio values as a function of prestrain for the X52 material for each charging method: a Electrolytic charging and b H 2 S Method. CLR – Crack Length Ratio CTR – Crack Thickness Ratio and CSR – Crack Size Ratio. ..................................................................................................... 59 Figure 4.20 Critical Crack Ratio values as a function of prestrain for the X60 material for each charging method: a Electrolytic charging and b H 2 S Method. CLR – Crack Length Ratio CTR – Crack Thickness Ratio and CSR – Crack Size Ratio. ..................................................................................................... 60 Figure 4.21 Critical Crack Ratio values as a function of prestrain for the X70 material for each charging method: a Electrolytic charging and b H 2 S Method. CLR – Crack Length Ratio CTR – Crack Thickness Ratio and CSR – Crack Size Ratio. ..................................................................................................... 61 Figure 4.22 Critical Crack Ratio values as a function of prestrain for the 100XF material for each charging method: a Electrolytic charging and b H 2 S Method. CLR – Crack Length Ratio CTR – Crack Thickness Ratio and CSR – Crack Size Ratio. .................................................................................... 61 Figure 4.23 a Crack length ratio dependence on tensile strength for the 0 pct prestrain condition: Electrolytic method EC and NACE Standard method H 2 S. b Crack thickness ratio dependence on tensile strength for the 0 pct prestrain condition. The vertical dashed lines represent the suggested 116 ksi threshold value for hydrogen assisted-cracking phenomena 47 and 48. The two horizontal dashed lines represent sour service requirements outlined by ISO 3183 52. ................................................. 62 Figure 4.24 a Crack length ratio and b crack thickness ratio dependence on hardness for the materials and pre- strain conditions evaluated by EC. The vertical dashed lines represent the suggested 22 HRC 248 HV threshold value for HIC susceptibility 47 and 48. The two horizontal dashed lines represent sour service requirements outlined by ISO 3183 52. ......................................................................... 63

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xi Figure 4.25 a Hydrogen concentration in as-received alloys. b Hydrogen concentration for each plate material as a function of prestrain. ...................................................................................................................... 64 Figure 4.26 Trapped hydrogen values after degassing of material. Materials were hydrogen charged with the EC method degassed at 50 °C for 72 hours sectioned and evaluated using the LECO® hydrogen analyzer to determine the trapped hydrogen content. ........................................................................... 65 Figure 4.27 Diffusible hydrogen values determined by two different evaluation methods on the four materials and prestrain conditions investigated: a Mercury displacement and b LECO® analysis. LECO® diffusible hydrogen values were determined using Equation 4.2. ........................................................ 67 Figure 4.28 Observation of hydrogen bubbles forming preferentially on the lower bottom end of the steel specimens during EC experiments. Sample numbers are shown on the steel sample in the image. ..... 68 Figure 4.29 Dependence of location of sample sample number on the amount of diffusible hydrogen measured by the mercury displacement method. Sample 1 represents the sample furthest away from the electrode connection. Dimensions of the sample were the full thickness of the plate 100 XF and X70 - 12.7 X60 - 9.5 and X52 - 19 mm x 20 ± 3 mm width x 25 ± 2 mm length. ............................... 69 Figure 4.30 Dependence of calculated Crack Length Ratio on the location of the examined faces for each material and prestrain condition for EC experiments a X52 b X60 c X70 and d 100XF. ..................... 70 Figure 4.31 HIC CLR dependence on a LECO® diffusible and b trapped hydrogen contents from EC. The two horizontal dashed lines represent sour service requirements outlined by ISO 3183 CLR 0.15 52. ...................................................................................................................................................... 71 Figure 4.32 Light optical macrographs taken on as-polished full width and thickness EC test specimens. Images are representative of the cracking behavior observed for materials at all prestrain levels. a X52 0 pct prestrain Face 1 b X60 3 pct prestrain Face 2 c X70 5 pct prestrain Face 1 and d 100XF 0 pct prestrain Face 2. .................................................................................................................................... 72 Figure 4.33 Secondary electron micrograph taken on the FESEM of the X52 EC 18 prestrain Face 1 condition. Etched with 2 pct nital. Evidence of transgranular 1and 3 and intergranular crack propagation 2.. 73 Figure 4.34 EDS mapping of non-etched X52 EC 18 prestrain Face 1condition. Primary crack propagation in close proximity to Al-O-Ca type inclusion. SEM image produced was taken in backscatter mode on the FESEM see pdf version for color. ................................................................................................ 74 Figure 4.35 Secondary electron micrographs taken on the X60 EC 0 prestrain Face 6 condition. Image captured using the ESEM. Etched with 2 pct nital. Features 1 and 2 show HIC around inclusions in the microstructure. ...................................................................................................................................... 75 Figure 4.36 Secondary electron micrograph taken on the X70 EC 5 prestrain Face 1 condition. Image shows the presence of an inclusion in the primary crack area. Image was captured using the ESEM. Etched with 2 pct nital. ............................................................................................................................................. 75 Figure 4.37 EDS mapping of non-etched X60 EC 0 prestrain Face 6 condition. SEM image was taken in backscatter mode on the FESEM see pdf version for color. The SEM image and EDS maps show the presence of MnS type inclusions in and distributed around the primary crack region. ................... 76

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xii Figure 4.38 EDS mapping of non-etched X70 EC 5 prestrain Face 1 condition. SEM image was taken in backscatter mode on the FESEM see pdf version for color. The SEM image and EDS maps show the presence of MnS type inclusions in and distributed around the primary crack region. ........................ 76 Figure 4.39 Secondary electron images take on a 100XF EC 0 prestrain Face 2 feature 1 shows transgranular cracking of acicular ferrite b 100XF EC 0 prestrain Face 2 feature 2 shows crack propagation across secondary microconstituent and c 100XF EC 2 prestrain Face 1 feature 3 shows nonmetallic inclusion. Image was captured using the FESEM. Etched with 2 pct nital. ..................... 77 Figure 4.40 EDS mapping of non-etched 100XF EC 0 prestrain Face 2 condition. SEM image was taken in backscatter mode on the FESEM see pdf version for color. SEM image reveals an globular shaped inclusion found outside the primary crack region. EDS mapping identifies the inclusion being of the mixed composition of Al-O-Ca-S. ........................................................................................................ 78 Figure 4.41 EDS mapping of etched 100XF EC 2 prestrain Face 1 condition. SEM image was taken in backscatter mode on the FESEM. Etched with 2 pct nital see pdf version for color. SEM image reveals an elongated inclusion found within the primary crack region. EDS mapping identifies the inclusion being of the mixed composition of Al-Mg-O. ....................................................................... 78

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xiii LIST OF TABLES Table 2.1 Chemical Composition of the Materials used in 41 wt pct ..................................................................... 16 Table 2.2 NACE Standard Test Results and Results Using Electrochemical Charging on X70 Plate 7 ................... 16 Table 3.1 Chemical Composition of As-Received Plate Steels in wt pct .................................................................... 18 Table 3.2 Mechanical Properties in the Transverse Direction of As-Received Plate Steels ........................................ 18 Table 3.3 Selected Prestrain Amounts for the Plate Steels .......................................................................................... 21 Table 3.4 Test Specimen Geometry for the X52 Specimens used in each Hydrogen Exposure Method .................... 23 Table 3.5 Test Specimen Geometry for the X60 Specimens used in each Hydrogen Exposure Method .................... 24 Table 3.6 Test Specimen Geometry for the X70 Specimens used in each Hydrogen Exposure Method .................... 25 Table 3.7 Test Specimen Geometry for the 100XF Specimens used in each Hydrogen Exposure Method ................ 26 Table 3.8 Design Issues and Solutions taken for the Electrolytic Charging Methodology .......................................... 27 Table 3.9 Measured values from Manual Image J Measurements of length ‘a’ and thickness ‘b’ of Cracks in Figure 3.11. The Calculated Crack Ratio Values are also shown. ……………………………………………34 Table 3.10 Uncertainty from the Method used to Measure Variables for use in Crack Ratio Calculation .................. 35 Table 3.12 Uncertainty Values for the Instruments Used to Calculate Diffusible Hydrogen ...................................... 42 Table 3.13 Total Uncertainty Calculated using Equation 3.9 for the Mercury Displacement Method ........................ 42 Table 4.1 Microhardness Data Taken for each Material and Prestrain Condition ....................................................... 54 Table 4.2 Experimental Matrix for the Characterization Study Conducted on the 100XF alloy ................................. 54 Table 4.3 Charging conditions that produced cracking in 100XF that was cathodically charged for 24 hours ........... 57 Table 4.4 Average crack depth in 100XF specimens at the current densities explored during EC experiments run for 24 hours: Extremes represent the minimum and maximum crack depth observed ………………58 Table 4.5 Average crack depth in 100XF specimens as a function of test duration during EC experiments run at 15 mA/cm 2 : Extremes represent the minimum and maximum crack depth observed ………………58 Table 4.6 Hydrogen trap sites found in iron table modified from 54 ....................................................................... 66 Table 4.7 Crack Depth Observed on all Sectioned Faces for each alloy after EC ....................................................... 72 Table 4.8 Summarized Cracking Behavior .................................................................................................................. 78

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xiv Acknowledgements I would like to thank the sponsors of the Advanced Steel Processing and Products Research Center for their support guidance and encouragement. Particular thanks are given to Evraz for allowing me to visit their RND facility to conduct the NACE standard test for my thesis work Thank you Laurie Collins and Shahrooz Nafisi for your support and guidance during my thesis. Thank you Nucor for also allowing me to visit and give you sample to conduct AFA analysis on Thank you Dan Edelman for all your support and guidance. Dr. Martins thank you for all your knowledge and guidance in this project without your expertise and communication the test cell would be nowhere close to where it is today. Thank you Lee and Jim Johnson in aiding me in the rebuild of the Bridgeport Knee mill as well as your knowledge you bestowed upon me. Additional thanks go to Matt Merwin U.S. Steel D. Bai and Rick Bodnar SSAB. I am grateful for my advisor Kip O Findley for supporting me through this program and increasing my knowledge and skills within the field of metallurgical and materials engineering. Finally I would like to thank my thesis committee for the knowledge and advice they gave to me during my thesis.

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1 CHAPTER 1: Introduction Plate steels used in pipeline applications are exposed to specific environmental challenges. One deleterious environment is the transportation of “sour gas” commodities. The presence of hydrogen sulfide gas promotes hydrogen-assisted cracking specifically sulfur inhibits the recombination of hydrogen into hydrogen gas thus facilitating hydrogen diffusion into steel. Steel processing has evolved over the years to produce steel products that exhibit higher resistance to hydrogen induced damage. Specifically processing strategies have included lowering sulfur contents higher cleanliness of steel produced 1 minimizing hydrogen contents controlled cooling and low moisture in raw materials 1 controlling the formation of specific microstructural constituents more homogenous 2 controlling nonmetallic inclusion shape control limiting the degree of elongation 6 and decreasing corrosion potentials while lower the diffusivity of hydrogen into the steel matrix alloying 6. Also casting techniques ingot versus continuous cast 4 deoxidation processes fully versus semi killed steels 6 rolling upon hot slab reduction angle of rolls 1 rolling 5 and controlled cooling minimization of ‘hard’ bands and alloy segregation 1 have been shown to affect hydrogen induced cracking resistance. Furthermore plate manufacturing results in bending strains in the material which may affect resistance to hydrogen induced damage through local increases in strength and decreases in ductility as well as increases in dislocation density and imposed residual stresses. The industrial standard for the evaluation of hydrogen induced cracking HIC is the NACE TM0284 standard. This method relies on naturally occurring corrosion reduction reactions at the interface between one of two different solutions A or B containing dissolved hydrogen-sulfide gas and the iron-containing base-metal. Thus hydrogen adsorption at the metal/solution interface and subsequent diffusion into the metal is accomplished. The reduction of the dissolved hydrogen sulfide occurs at cathodic sites on the metal surface and the adsorption of atomic hydrogen is facilitated by the accompanying sulfide anion reaction-product that functions as a hydrogen association-inhibitor reduces the rate of formation of diatomic-hydrogen gas. However there is inherent danger in handling hydrogen sulfide gas and special procedures are required to use it for hydrogen-assisted cracking evaluation purposes. 1.1 Research Objectives The scope of the current project is to develop a reproducible methodology using electrolytic charging as a means to introduce hydrogen into steels to assess their resistance to hydrogen-assisted cracking. Electrolytic charging has been commonly used to pre-charge steel specimens with hydrogen before mechanical testing or to charge hydrogen into specimens while applying a load however it has not been extensively employed to assess hydrogen-induced cracking resistance 6. Electrolytic charging at prolonged test durations was explored by Rosner 7 who found that electrolytic charging produced similar results to the NACE Standard test mentioned above a promising initial result.

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2 The current project utilizes four low carbon plate steel grades X52 X60 X70 and 100XF to evaluate hydrogen susceptibility in relation to: 1. As-received mechanical properties i.e. strength hardness and ductility 2. The introduction of cold work in the form of uniaxial prestrain 3. Diffusible hydrogen amounts along with analysis of trapped hydrogen after degassing 4. Parent microstructure along with second phase microconstituents and 5. Presence of nonmetallic inclusions i.e. type morphology and/or size.

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3 CHAPTER 2: Background In the following section the various proposed mechanisms for hydrogen-assisted cracking are outlined and discussed. Variables such as selection of raw materials for steelmaking fabrication processes that occur during the production of steel plates and the final forming methods are presented in relation to hydrogen induced damage. Two experimental methods to introduce hydrogen into steel to study the effects of hydrogen are also discussed. 2.1 Hydrogen Entry into Steel and Hydrogen-Assisted Cracking Mechanisms Hydrogen can be introduced into steel though steelmaking processes and in-service environmental exposure. Exposure to hydrogen from steelmaking comes from melting of the raw materials casting of ingots or slabs and subsequent processing techniques e.g. electro-plating and pickling. In-service environmental hydrogen exposure has increased due to the increasing demand for oil and gas commodities leading to the exploring and harvesting of reserves that contain hydrogen sulfide gas and/or sulfur as impurities. In the oil and gas industry the terms ‘sour’ and ‘sweet’ are used to denote the concentrations of the two impurities found in the commodity. For gas commodities a ‘sour’ gas is defined as when the concentration of hydrogen sulfide gas exceeds 5.7 milligrams per cubic meter of natural gas. Oil wells are considered ‘sour’ when sulfur levels are at concentrations that exceed 0.5 wt pct. At lower levels of hydrogen sulfide gas or sulfur the gas or oil commodity is described as ‘sweet.’ The production interaction and subsequent damage from hydrogen in steels are schematically shown in Figure 2.1. The production of nascent hydrogen atoms is generated through the naturally occurring corrosion reaction → 2.1 in the sour environment 5. Hydrogen ions are attracted to the metal interface due to the presence of excess electrons. Once at the interface the nascent hydrogen can either recombine into hydrogen gas and bubble off the surface or diffuse into the steel and cause internal cracking and/or blistering at the surface. The diffusion of hydrogen into the steel is assisted by the presence of sulfide ions. The sulfide ions slow the recombination reaction of hydrogen ions to form hydrogen gas within the system. Internal cracking is generated though interactions hydrogen has with specific features such as inclusions in the metal matrix. Hydrogen-assisted cracking phenomena are caused by the following proposed mechanisms 8: internal hydrogen pressure blister or void formation surface adsorption decohesion and enhanced plastic flow. Surface adsorption affects the energy needed decreases to form brittle cracking 8. Decohesion occurs because the dissolved hydrogen weakens interatomic bonds therefore lowering the ductility of a material. These two mechanisms surface adsorption and decohesion require an applied stress either external or residual to observe their consequences. However the other two mechanisms internal hydrogen pressure and enhanced plastic flow can occur without an applied external stress they are related to an internal stress caused by increased levels of hydrogen in specific areas. These mechanisms lead to premature failures or costly repairs of the steels used in service.

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4 Figure 2.1 Schematic representation of the dissociation of hydrogen sulfide gas at the metal/solution interface and subsequent diffusion into the metal to areas where hydrogen-assisted cracking can occur 5. Once hydrogen is introduced into the steel it can lead to different forms of cracking such as sulfide-stress- cracking SSC stepwise cracking SWC hydrogen induced cracking HIC hydrogen blister cracking and stress- oriented-hydrogen-induced-cracking SOHIC. SSC is a form of hydrogen-assisted cracking that occurs in high strength steels 100 ksi yield strength or in areas of localized hardness i.e. ‘hard bands’ produced though alloy segregation or heat affected zones found close to weldments. This type of cracking is produced through exposure to hydrogen while under tensile stress. SSC is generally intergranular and is primarily oriented perpendicular to the applied tensile stress 9. SOHIC is a form of hydrogen-assisted cracking where the damage caused through the interaction with hydrogen is indicative of an applied stress. SOHIC is characterized by the interlinking of microscopic cracks both on the surface and within the steel that are oriented perpendicular to the applied stress as well as the plane defined by nonmetallic inclusions. Areas that are highly susceptible to SOHIC are the same as those described for SSC and material adjacent to the weld seam in pipeline steels. HIC covers a broad range of cracking behavior that occurs through the interaction with hydrogen. Within HIC terms such as SWC hydrogen pressure cracking and hydrogen blister cracking are interchanged frequently. For the purposes of this document all will be combined and referred to as HIC. HIC is caused through the internal interaction of hydrogen with the steel in the absence of an externally-applied stress. Hydrogen is believed to affect a material through its interaction with crystalline features such as dislocations and/or secondary micro-constituents such as carbides and non-metallic inclusions 8. At these features hydrogen becomes trapped and builds to pressures greater than the local strength of material causing internal micro-cracks to form. Enhanced plastic flow has been one of the more accepted mechanisms associated with HIC and is shown schematically in Figure 2.2. This proposed mechanism for enhanced plastic flow is as follows: 1 hydrogen diffuses through the metal matrix and becomes trapped at the metal/inclusion interface 2 the presence of molecular hydrogen causes a separation between matrix metal and inclusions 3 plastic deformation occurs around the crack tips 4 plastic regions are embrittled by hydrogen stage 1 and 5 cracks propagate through the embrittled regions in a direction normal to blisters cracks stage 2 3. Scenarios A B and C in figure 2.2 represent various configurations of inclusion size amount and spacing. Scenario A represents two elongated type inclusions that are separated B represents a cluster of smaller inclusions that within close proximity to one another and C shows a very

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5 elongated inclusion in the presence of two minor inclusions. The crack morphology is altered from scenario A to C due the differences in the features of the inclusions. Figure 2.2 Schematic representation of enhanced plastic flow mechanism for hydrogen embrittlement. Parts A B and C show various configurations of inclusion size/elongation distribution and amount and the resulting morphology of cracks 3. 2.2 Metallurgical Variables This section evaluates important metallurgical variables that are considered when producing plate steels. Variables such as different casting techniques microstructure nonmetallic inclusions and alloy chemistry are important to monitor to increase steel resistance to HIC. Additionally the effect of thermomechanical processing route and forming steps required for final pipe production from plate product are also considered. Before addressing these variables it is important to distinguish the different classes of pipeline steels that are used within the oil and gas industry. These steels are as follows: 1 conventional steels 2 low sulfur conventional steels 3 “HIC-Resistant” Steels and 4 ultra-low sulfur advanced steels 10. The term “conventional steel” refers to commercially produced hot rolled or normalized steel. When the environment is not challenging the steels generally have a moderate level of inclusions and higher amounts compared to others mentioned here of sulfur present in the final chemistry of the steel ≥ 0.010 wt pct sulfur 10. These steel grades when placed into sour environments exhibit high degrees of hydrogen induced damage which is to be expected from the higher levels of sulfur and inclusions present in the steel. Low sulfur conventional steels are produced in a similar manner to conventional steels but with more controlled sulfur levels. The sulfur levels range from 0.003 to 0.010 wt pct. Due to the decreased levels of sulfur the amount of inclusions present is lower and these steels exhibit improved mechanical properties compared to conventional steels. When placed into a sour environment low sulfur conventional steels perform better than conventional steels but may still exhibit a high susceptibility to cracking at moderate to severe service conditions 10.

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6 HIC-Resistant steels denote steel that has been engineered at multiple levels of processing such that this material can be placed into moderate and severe environments and not fail due to hydrogen induced damage. Processing during ladle metallurgy involves the reduction of sulfur to levels designated as ultra-low i.e. ≤ 0.002 wt pct and a possible addition of calcium to aid in nonmetallic inclusion shape control during subsequent processing 10. After hot rolling these steels undergo a normalizing heat treatment to modify the rolled microstructure to improve the resistance to hydrogen induced damage. Ultra-low sulfur advanced steels have been specifically engineered to resist HIC under severe service conditions and improve resistance to SOHIC through the use of modern steelmaking and processing techniques. These steels like the HIC-Resistant steel have ultra-low sulfur contents and lower carbon equivalents at similar tensile strengths as conventional steels such as ASTM A516-70. Ultra-low sulfur advanced steels undergo thermo- mechanically controlled processing TMCP and/or accelerated cooling techniques to produce ferritic/bainitic microstructures where banding is minimal or not present 10. Because of the TMCP attention to sulfur contents and nonmetallic inclusion shape control these steels exhibit the highest resistance to hydrogen induced damage of the four grades of steel mentioned. 2.2.1 Deoxidation and Casting Techniques Resistance to hydrogen induced damage is altered by casting techniques due to the control over alloy segregation and regions of nonhomogeneous microstructure. Variables such as deoxidation process ‘fully’ versus ‘semi’ killed casting technique ingot versus continuous cast super heat in the tundish casting speed and position of rolls can alter the how the final steel product solidifies and where nonmetallic inclusions are present. Two types of deoxidation processes can be used before casting. The terms “fully-killed” and “semi-killed” are in reference to the deoxidation process before casting and the elements that are added to achieve deoxidation. In “fully-killed” steels aluminum is added because its strong affinity for oxygen limits the gas evolution in the bath of liquid steel 11. Aluminum oxides form in the liquid and float to the top but some oxide inclusions remain in the steel upon casting. In comparison “semi-killed’ steels have ferrosilicon added to achieve deoxidation within the bath of liquid steel 12. “Semi-killed” steels produce heats that have more homogenous chemistries as well as more uniform mechanical properties after rolling operations 12. In a study conducted by Moore and Warga 12 on 14 candidate steels for pipeline HIC susceptibility was greater in fully-killed steels than semi-killed steels. The susceptibility was independent of the varying levels of sulfur manganese and silicon in the alloys but did depend on the amount of aluminum. The higher amount of deoxidation that occurs in the fully-killed steels leads to higher solubility of sulfur in the melt. The increase in solubility results in more Type I ellipsoidal manganese sulfides MnS upon solidification which are subsequently elongated during rolling leading to Type II stringer-like MnS 12. The Type II MnS inclusions increase the degree of cracking in plate steels subjected to sour environments. In the semi-killed steels the sulfides present after solidification and rolling operations are more globular increasing the resistance to hydrogen-assisted cracking phenomena. Herbslab et al. 4 showed that continuous casting should be used instead of ingot casting because the steelmaker has better control over the position of the nonmetallic inclusions. Nieto et al. showed that continuous casting enabled them to control the final solidification of slabs to minimize centerline segregation. They also showed

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7 that continuous casting low superheats in the tundish 250°C and slower casting speeds produced plate steels API X65 with high resistance to HIC 1 due to no centerline segregation homogeneous chemistry though thickness and microstructural uniformity. In contrast when plates are produced from large ingots it has been shown that HIC susceptibility scales with the position in the ingot 3. In a study conducted by Ikeda et al. 13 samples were taken from four different locations within a large ingot and tested for HIC susceptibility. The HIC susceptibility in relation to these regions is shown in the ingot schematic in Figure 2.3. The positions examined were1/6 of the width W 1/3 W 5/12 W and 11/12W from the surface as well as the lower section 1/6 of the height of the ingot away from the slag layer Figure 2.3. This study showed that in the area labeled as ‘A’ the outer edge and lower portion of the ingot the HIC susceptibility of the final plate is not severe. Towards the center of the ingot i.e. positions of higher concentrations of impurities denoted as areas ‘B and C’ the HIC susceptibility increases moderately compared to ‘A’. Samples taken from the inner third of the ingot section D produced the lowest resistance to HIC due to the large amounts of alloy segregation and high concentrations of nonmetallic inclusions 13. The importance of this study showed that steels produced from ingots have varying degrees of susceptibility and those from sections B C and D should not be used in sour environments. Figure 2.3 Schematic representation of the influence of ingot material location on HIC susceptibility of large sized ingots used to produce hot-rolled product 3.

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8 2.2.2 Microstructural Variables This section highlights the different microstructural variables that are present in the final plate steel and how they influence the susceptibility to hydrogen-assisted cracking. The parent microstructure presence of secondary microconstituents non-homogenous microstructure chemical make-up of nonmetallic inclusions size and shape of nonmetallic inclusions and finally alloying influence susceptibility to hydrogen-assisted cracking. In general it is recommended by the National Association of Corrosion Engineers NACE that plate steels for use in sour service environments should not exceed a hardness value of 22 – 25 Rockwell C HRC 3. Based on this recommendation steels with an ultimate tensile strength of 116 ksi 800 MPa or greater should not be considered for use in hydrogen environments. It is accepted in literature 14 and 15 that as the strength of the material increases the tolerance to hydrogen concentration decreases and its overall mechanical properties toughness fracture toughness tensile strength are altered by the interaction with hydrogen. “Harder” microstructures inherently are generally less ductile and require lower concentrations of hydrogen to cause severe damage to the steel. Acicular ferrite grains have a three-dimensional morphology of thin lenticular plates where grain size is characterized by an aspect ratio that ranges from 3:1 to 10:1 16. Acicular ferrite steels exhibit high strength coupled with high ductility and have a higher resistance to hydrogen compared to polygonal ferrite or mixed ferrite/pearlite microstructures. The degree of brittle intergranular fracture from the presence of hydrogen is reduced in acicular ferrite compared to polygonal ferrite 14 and 15. In one study tempering acicular-ferritic steels at 650 °C for 30 minutes increased the resistance to HIC thorough the elimination of hard transformation products 3. The tempering step allowed sufficient carbon diffusion out of martensitic islands decreasing the susceptibly of cracking around these features. Increases in HIC resistance were also tied to increases in the uniformity of the microstructure increases in fracture toughness lower hardness and internal stress relief. Grain size effects have been shown to alter HIC susceptibility in polygonal ferritic steels. Grain coarsening in polygonal ferrite has pronounced adverse effects on hydrogen-assisted cracking due to the ease of transgranular crack propagation in larger grains. Grain boundaries have been shown to reduce crack propagation through crack tip shielding 17. Therefore increasing the amount grain boundaries and decreasing grain size increases the degree to which a crack interacts with grain boundaries reducing hydrogen-assisted cracking. However at very fine grain sizes it has been hypothesized that the increased grain boundary surface area associated with small-grained structures increases the probability for monoatomic hydrogen to segregate to grain boundaries increasing the susceptibility to hydrogen-assisted cracking 18. Pearlite in steel interacts differently with hydrogen than ferrite. Cementite in the pearlite structure acts as a retardant for hydrogen diffusion through the microstructure but hydrogen reaches similar concentration levels to purely ferritic steels after longer exposure times 19 and 20. It has been reported that HIC susceptibility scales with the degree of banding present in the parent microstructure 3 10 and 21. Hydrogen diffusion through a mixed ferrite/pearlite microstructure is shown Figure 2.4. The arrows indicate the path by which hydrogen diffuses though microstructure. Figure 2.4a shows a lower degree of banding compared to figure 2.4b. As hydrogen moves through the microstructure pearlitic colonies hinder its diffusion and thus less banded steels inhibit hydrogen diffusion.

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9 Thus it was observed that an increase in the banding Figure 2.4b increased the amount of area affected by cracks 2. Hydrogen is irreversibly trapped at the interface between the ferrite and cementite phases that make up lamellar pearlite. After long exposure times the interface of the cementite becomes saturated by hydrogen causing cracking. a b Figure 2.4 Schematic diagram illustrating the hydrogen penetration through two varying degrees of banding in a ferrite/pearlite microstructure: a lower degree of banding and b higher degree of banding. In a the hydrogen penetration is hindered by pearlite along the ferrite regions. Adapted from 19. Taira et al. 2 showed that by quenching at 900 °C and tempering at 650 °C QT controlled rolled steels there was a reduction in the amount of cracking compared to the same material that was only controlled rolled. The increased resistance was attributed to the elimination of the pearlitic structure and increased uniformity of the microstructure compared to controlled rolled steels of the same grade. The quenching and tempering process produced tempered bainite resulting in an overall increase resistance to HIC. Other studies 22 and 23 on QT steels have also reported an increase in the resistance to HIC. It becomes apparent that for pipeline steels that are subjected to sour/severe environments that the control of nonmetallic inclusions is paramount to ensure resistance to hydrogen-assisted cracking. Inclusions that form in the final plate product are dependent on liquid metal processing practices. Practices such as stirring how long when how vigorous timing of alloy additions shrouding gas layers above the steel controlling slag and many others have a very strong effect on inclusion counts in the final steel 24. The region of discontinuity around the inclusion and matrix creates a trap site for hydrogen accumulation favoring crack nucleation and propagation in those regions. The length of the inclusion is of greater importance than its thickness 15 with respect to the rolling direction or parallel to the banding direction. The interface area along the length of the inclusion is greater making the material more susceptible to brittle fracture normal to these interfaces. Figure 2.5 is a schematic representation of examples of inclusions that might be present in fully killed steels 25 after casting and rolling operations. Alumina inclusions during casting form a dendritic structure but upon rolling operations are broken up and distributed with respect to the rolling direction. At higher calcium oxide to alumina ratios the inclusion tends to become more globular in shape and will have minimal shape change after rolling operations. In contrast MnS inclusions elongate greatly during rolling which is very detrimental to resistance to hydrogen assisted cracking 3 5 21 and 26. During casting MnS inclusions are globular but the

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10 inclusion becomes elongated after rolling operations. This elongation allows for an increase in surface area for hydrogen adsorption along a single plane increasing the degree of cracking observed in the microstructure 3 26. Because nonmetallic inclusions cannot be eliminated the final example found in Figure 2.5 is an example of the type of inclusion that would be more resistant to HIC. The 12 CaO • 7 Al 2 O 3 inclusion forms first during casting and then the sulfide ring CaS/MnS forms around it. The sulfide ring allows for better control of inclusion shape as well as size and distribution within the metal matrix. Controlling these variables with additions of calcium aids in enhancing resistance to HIC. Figure 2.5 Examples of nonmetallic inclusions that form in fully killed steels. Shown are the changes of morphology due to rolling operations after casting. Adapted from 25. Alloying elements and impurities may affect interactions of steel with hydrogen by changing the corrosion potential of the steel poisoning surfaces such as grain boundaries forming precipitates that act as traps and forming protective layers 3 6 26 and 27. It is shown in literature that in environments where protective surface films can form i.e. not severe sour service copper additions greater than 0.20 wt pct increase plate steels resistance to HIC 3 21 27 – 29. However when copper is present with molybdenum nickel or tungsten in the steel it has been documented that the resistance to HIC is decreased due to their interactions increasing corrosion rates along with hydrogen adsorption 22 and 30. In chromium nickel and cobalt-bearing alloys the addition of these elements lowers hydrogen adsorption during exposure thus increasing the resistance to HIC. Certain irreversible traps precipitates have been shown to increase the resistance to hydrogen-assisted crack phenomena 19 and 31. The overall effectiveness of precipitates to increase resistance to HIC is based on the size and the distribution of that precipitate in the matrix 19 and 31. Precipitates have high enough binding energy i.e. 60 kJ/mole to act as irreversible traps and when they are distributed homogenously small compared to inclusions and present in sufficient quantities they increase the resistance to HIC 31. When the precipitates interact with hydrogen they minimize the amount of hydrogen that can diffuse elsewhere into the material and initiate cracks around other irreversible traps such as inclusions. Assuming the precipitates are distributed

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11 homogenously and in sufficient quantities they can minimize the amount of cracking that occurs by allowing hydrogen to be more evenly distributed through the microstructure instead of at areas of higher crack susceptibility. 2.2.3 Mechanical Processing Effects During the forming of plate steels from continuous cast slabs or ingots parameters that may alter the resistance to HIC are start temperature for accelerated cooling cooling rate and finishing temperature 3 5 and 20. Strain introduced through cold rolling operations has been reported to have both detrimental and beneficial effects on plate steel performance in hydrogen environments. Further processing of plate into seam-welded pipe introduces plastic deformation and residual stresses into the final product. This section highlights the important variables to consider when rolling cooling and forming plate steels into pipeline steel products. During hot rolling of plate steels at low finishing temperatures residual stresses are generated within the plate steel and subsequent hot working does not relieve these stresses. Finishing temperatures of 900 °C produced steels that had higher resistance to HIC when compared to the same steels rolled at a low finishing temperature of 790 °C because the higher finishing temperature allowed for higher amounts of residual stresses to be relieved 19. Moore and Warga 12 showed that with decreasing finishing temperature the elongation of Type I MnS inclusions increased. This increase in elongated MnS within the plate steel increased the overall susceptibly to HIC. In a study conducted by Shinohara and Hara et al. 5 when the start temperature of accelerated cooling was held above approximately 10 °C of the A r3 temperature HIC susceptibly was lowered. Similarly they found that when the cooling rate of the steel is below 10 °C/s the banding present in their X70 pipeline steel at mid-thickness was still present. The banding present at mid-thickness created areas of preferential cracking increasing the susceptibility to HIC. A good approximation for the levels of residual stresses is the X-ray diffraction line broadening 3 and 20. Line broadening also relates to increases in increased dislocation density. Figure 2.6 shows the line broadening parameter and crack length ratio extent of cracking in a low carbon 0.07 wt pct micro-alloyed Nb and V plate steel. The HIC test used to evaluate this material was the NACE standard test TM0284. Crack length ratio the definition of which is provided in section 3.5 was calculated as per the standard upon optical microscopy inspection of sectioned faces. The HIC resistance decreases while the line broadening parameter increases as the amount of strain in the material increases. Cold reduction from 0 to 10 pct reduces the crack length ratio from a value of 60 pct to 20 pct. The largest change in susceptibility is between the values of 0 to 5 pct cold rolling. This is important to note because beneficial effects of cold rolling are observed at relatively low amounts of cold reduction i.e. strain. The uniformity of strain through the thickness at these low reductions may also play in a role in HIC resistance. The effect of larger amounts of cold reduction on HIC susceptibly is shown in Figure 2.7. Three steels A B and C had similar levels of cold reduction introduced and were exposed to the NACE standard test TM0284. Compositions for the steels are as follows: Steel A - 0.04-C 1.04-Mn 0.004-S Steel B – 0.05-C 1.24-Mn 0.005-S Steel C – 0.09-C 1.01-Mn 0.008-S all values are in wt pct all 3 steels were micro-alloyed Nb and V plate steels. The plate thickness was 9.52 mm for A 19 mm for B and 14.3 mm for C. The experimental methodologies for cold rolling operations such as strain per pass were not described. Steels A and C were produced for sour service

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12 applications and steel B was a low sulfur grade steel. Steel C had the lowest degree of HIC at the highest cold reduction. This was attributed to it being continuously cast compared to steel A and B that were ingot cast. The increase in non-homogenous microstructure in steel A increased its susceptibility at amounts higher cold reduction compared to the other alloys. The overall trend observed for the three steels is very similar and can be described as an S-curve Figure 2.7. At values of cold reduction greater than 50 pct the crack length ratio reaches saturation and thus increasing the value of cold reduction to greater values has minimal effect to increase the degree of cracking. It is interesting to compare the results from Figure 2.6 and 2.7 where relatively small amounts of cold reduction have a beneficial effect on HIC resistance in Figure 2.6 while much larger reductions have a detrimental effect. Other studies 32 and 33 have reported that values of cold strain from 2 to 16 pct increase susceptibility to HIC. The increase in resistance at low strains was attributed to decreases in the permeability of hydrogen into the metal matrix and a more uniform distribution of hydrogen in the matrix 20. When hydrogen is unevenly distributed through the metal matrix areas susceptible to hydrogen such as nonmetallic inclusions reach levels of hydrogen in sufficient amount to initiate HIC. Whereas if hydrogen is more evenly distributed through the metal matrix due to an increased level of traps created by cold reduction HIC susceptibility decreases 20. Figure 2.6 HIC susceptibility crack length ratio ratio as a function of cold rolling reduction. Lattice strain was determined through X-ray diffraction after cold rolling operations denoted as line broadening in the above figure. HIC testing was performed with respect to the NACE standard test TM0284. Steel used for HIC testing was a 0.07-C 1.22-Mn 0.006-S all values in wt pct micro-alloyed with Nb and V heavily controlled rolled steel. Adapted from 20.

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13 Figure 2.7 Effects of cold reduction at levels greater than 10 pct. Steels A and B have similar carbon equivalents 0.31 and were produced from an ingot whereas steel C was produced by continuous casting and has a higher carbon equivalent of 0.39. Steels A and C would be categorized are HIC- Resistant steels and steel B is a low sulfur grade steel. See text for more detailed chemical compositions of Steels A B and C. Adapted from 20. 2.3 Experimental Methods to Test Plate Steels for HIC Susceptibility In this section two experimental methods to evaluate pipeline steel susceptibility to HIC are outlined and discussed. The industrial standard developed by NACE involves the use of hydrogen sulfide gas to achieve hydrogen entry into the steel matrix. Electrolytic charging uses electrochemical reactions and polarization in an aqueous solution to introduce hydrogen to achieve similar effects as produced by the NACE standard test. 2.3.1 NACE Standard Test TM0284 For the introduction of hydrogen into the pipeline steels the NACE standard 34 method uses naturally occurring corrosion reactions within an aqueous system saturated with hydrogen sulfide gas. These principals are schematically shown in Figure 2.8. At the metal/solution interface anodic sites i.e. elemental iron are oxidized through the following equation: 2.2 while the hydrogen sulfide is oxidized according to

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14 ↔ 2.3 due to the presence of water in the solution. Hydrogen is then reduced according to the following equation: 2.4 These reactions create an electrochemical cell of reactions within the system. The presence of sulfide ions in solution due to the oxidation of hydrogen sulfide gas is known to suppress the rate at which the reaction in Equation 2.4 occurs. Hydrogen ions are attracted to the metal/solution interface due to the presence of electrons at the surface from the oxidation of iron 35 and 36. The interaction of hydrogen at the surface allows for hydrogen to recombine into hydrogen gas and bubble off to the surface into solution Figure 2.8a or diffuse into the steel matrix and potentially cause damage in the form of HIC Figure 2.8b. Figure 2.8 Schematic illustration of the two reactions for hydrogen that occur at the metal/solution interface. The reaction a occurs at a much higher rate than b. Reaction kinetics of a are suppressed in sour service applications due to the presence of sulfur ions Equation 2.3. Reaction b causes subsequent HIC within the material through hydrogen interactions with features in the metal matrix 36. The NACE Standard TM0284 34 is a standardized test method which was established to enable consistent evaluation of pipeline steels and their performance under sour environments. This performance is evaluated by the amount of HIC-induced damage the test generates in the specimen. The test is not designed to evaluate other adverse effects from a sour environment such as pitting weight loss from corrosion or sulfide stress cracking. The conditions of the test are not designed to simulate any specific pipeline or process operation 34. For the test unstressed specimens are exposed to one of two test solutions solution A – a sodium chloride acetic acid NaCl CH 3 COOH solution saturated with hydrogen sulfide gas at ambient temperature and pressure or solution B – a synthetic seawater solution saturated with hydrogen sulfide gas at ambient temperature and pressure. The specimens are subject to a 96 hour test in either solution after which they are removed and evaluated. Tests are conducted in an airtight vessel with adequate space for the test specimen along with provisions for gas purging and introduction of hydrogen sulfide gas into the vessel. The ratio of volume of test solution to the total exposed surface area of the test specimen is a minimum of 3 mL/cm 2 . The evaluation of HIC susceptibly is based on the calculation of three parameters crack length ratio CLR crack thickness ratio CTR and crack size ratio CSR. CLR is the sum of all the length of the cracks found on the evaluated faces normalized by the total width of the face that is examined. Higher values of CLR without inspection of microstructure could indicate higher degrees of cracking oriented with microstructural banding or elongation of inclusions. CTR is a measure of how much of the plate thickness is affected by HIC cracking. CTR values are a good indication of the linkage of cracks within the microstructure based on inclusion size and distribution as shown

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15 in Figure 2.2b. CSR is a measure of the total area of the cross-sectioned face that is affected by cracking. Because this sums the area of the cracks in reference to total cross-section high values of CSR would indicate large cracks were produced both in length and thickness of the crack. The definitions of CLR CTR and CSR are presented in greater detail in section 3.5. 2.3.2 Electrolytic Charging Electrochemical polarization ƞ is the measure of the potential change in a system relative to the equilibrium half-cell reactions the potential change is caused by changing the rate at which half-cell reactions occur at the interface of the metal/solution. Cathodic polarization causes an excess of electrons at the metal/solution interface and the surface potential becomes negative which slows the anodic dissolution reaction rate. Therefore by definition the value of ƞ for cathodic polarization is negative. For anodic polarization the surface potential is positive and the value of ƞ is positive. Another term used for cathodic polarization is overpotential 8. These surface potentials and their effects on reaction rates are related to the kinetics of the system not the thermodynamics. It should be noted that the reaction rates within the system do not correspond to the half-cell electrode potentials but rather to the current density associated with the reaction. The reactions are limited by kinetic surface reaction rates mass transfer from both the material and aqueous solution and potential concentration gradients found within the system 8. Electrolytic charging is conducted by selectively polarizing a two electrode system by using a DC power supply that has negative and positive output terminals. The steel specimen serves as the cathode and another material such as graphite serves as the anode. The galvanic couple between the steel specimen and anode material creates the cathodic polarization mentioned in the previous paragraph which allows hydrogen to be attracted to the steel specimen. Electrolytic charging is aided by the use of elemental additives such as arsenic or sulfuric species present in the electrolyte. The additives are introduced into the system as soluble compounds e.g. As 2 O 3 NaAsO 2 CS 2 which inhibit the formation of diatomic hydrogen from monatomic hydrogen 37. By doing so the uptake of hydrogen in the steel is enhanced increasing the efficiency of the charging setup. In literature the additives are referred to in a variety of terms such as hydrogenation promoter 37 and 38 hydrogen recombination inhibitor agents 38 or poison 38 - 40. The effectiveness of the inhibitors can depend on their relative amounts in the electrolyte 40. An electrolytic charging methodology has been commonly used to pre-charge steel specimens with hydrogen before mechanical testing or to charge hydrogen into specimens while applying a load. However it has not been extensively employed to assess HIC resistance 6. Perez Escobar et al. conducted a study that assessed the damage in high strength steels in relation to electrolytic charging conditions 41. While this study has not designed to assess HIC resistance it does provide some insight on charging conditions that might produce HIC damage. Four multi-phase high strength steels and pure iron were selected. The chemical compositions of the alloys used are displayed in Table 2.1. FB450 is a ferritic- bainitic mixed microstructure steel the TRIP700 is a multiphase steel containing ferrite bainite and retained austenite the DP600 is a dual phase ferrite-martensite steel and the S550MC is a high strength low alloy that has a mixed ferritie-pearlite microstructure with Ti-Nb precipitates 41. The main objective of this study was to identify electrolytic charging conditions that produced blisters on the surface of the materials investigated. To achieve this

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16 electrochemical variables such as current density applied test duration aqueous solution acidic or basic and presence of additives As 2 O 3 or thiourea CH 4 N 2 S were varied 41. Blister formation increased at both higher current densities and longer test durations. For acidic aqueous solutions at similar charging times and applied current densities the blister formation was greater. Pure iron was highly susceptible to blister formation in comparison to high strength ferrite bainite steel. In the absence of blisters on the surface Perez Escobar et al. sectioned the high strength steels and examined for internal cracks. For the FB450 TRIP700 and DP600 cracks were generated at the center of the sample around elongated MnS inclusions 41. Table 2.1 Chemical Composition of the Materials used in 41 wt pct Material/Element C Mn Si Other FB450 0.07 1.00 0.10 0.5 – 1.0 Cr TRIP700 0.17 1.60 0.40 1 – 2 Al 0.04 – 0.1 P DP600 0.07 1.50 0.25 0.4 – 0.8 Cr + Mo S550MC 0.07 0.95 0.0 0.08 – 0.12 Ti + Nb Pure Iron 0.0015 0.0003 0.0 0.02 Al P Most electrochemical charging times have been short i.e. less than 6 hours which does not allow for hydrogen to reach high enough levels of diffusible hydrogen to cause permanent internal damage in the form of HIC. In a recent ASPPRC study X70 plate steel was charged with hydrogen by electrochemical charging in a 1 normal sulfuric acid with 20 mg/L of As 2 O 3 additive at test durations up to 24 hours 7. The same X70 plate was also exposed to the NACE standard test TM0284 methodology. After the exposure to each method the specimens were sectioned and evaluated for internal cracks generated by hydrogen. At test durations of 24 hours in electrolytic charging the critical cracking parameters were on the same order of magnitude as the results produced by the NACE standard method conducted by U.S. Steel 7. Table 2.2 shows the results obtained from the NACE Standard test and the electrolytic testing methodology. In general the CTR CTR and CSR results from electrolytic testing range were greater than or equal to results obtained from the NACE standard test. Table 2.2 NACE Standard Test Results and Results Using Electrochemical Charging on X70 Plate 7 NACE standard test conducted by U.S. Steel B. Rosner Electrochemical Testing CSR CLR CTR Test No. CSR CLR CTR 0.1 6.0 0.3 1 0.33 24.7 2.9 1.5 16.7 3.1 2 0.05 9.40 1.60 3 0.33 15.50 3.40 4 0.06 6.65 1.25 5 0.10 2.95 3.60 6 1.46 41.5 10.8 7 1.88 43.4 15.1

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17 CHAPTER 3: Experimental Design and Methods This chapter presents the experimental methodology to create an electrolytic charging apparatus and procedure to evaluate plate steel susceptibility to hydrogen damage. 3.1 Experimental Design Rosner 7 explored electrolytic charging to develop a reproducible methodology to investigate hydrogen induced cracking and hydrogen damage in steels. In the study he showed that electrolytic charging could serve as a viable substitute to evaluate hydrogen susceptibility if there were modifications to increase the stability of methodology at prolonged test durations 7. These issues and solutions are presented and discussed in section 3.5.1. A newly designed test cell for electrolytic charging was fabricated. Characterization and validation of the electrolytic charging apparatus was undertaken. Alterations to certain parameters of the test apparatus and sample were conducted on a highly susceptible material 100XF to optimize the setup for HIC susceptibility experiments. Parameters such as current density voltage test duration and surface finish of cathode were identified as variables that could alter the reproducibility of results obtained from this methodology. The validation of the proposed experimental procedure was evaluated by comparing results from the industrial NACE Standard for HIC evaluation. The electrolytic charging apparatus was used to investigate HIC susceptibility of four different plate steel grades which were also subjected to cold-working before charging uniaxial prestrain. The alloys were all low carbon plate steels with a range of tensile strengths. Two of the alloys X52 and X60 were designated as HIC resistant grades. After exposure to the hydrogen environments either wet hydrogen sulfide gas or electrolytic charging through-thickness cross-sections were used to measure the extent of cracking. Further characterization of hydrogen damage was undertaken to explore effects of microstructure nonmetallic inclusions associated with cracks and diffusible and trapped hydrogen amounts. These methods are presented and discussed in subsequent sections of this chapter. 3.2 Experimental Materials and Methods 3.2.1 Materials The chemical compositions of the steel plates used in this study are shown in Table 3.1. The effects of prestraining were investigated using 100XF plate steel from the M.S. thesis of B. Farber 42 X70 plate steel from the M.S thesis of H.M. Al-Jabr 43 X60 plate steel received from EVRAZ North America and X52 plate steel received from SSAB. The X70 was provided by Essar Algoma Steel in 45 x 30 x 1.27 cm 18 x 12 x 0.50 in plates. The 100XF plate steel was received as 91 x 91 x 1.27 cm 36 x 36 x 0.5 in plates. The X60 plate steel was received as 84 x 34 x 0.95 cm 33 x 13.5 x 0.375 in plates. The X52 plate steel was received as 45 x 30 x 1.9 cm 18 x 12 x 0.75 in plates the plates came from the center of the plate stock. The X52 and X60 plate steels have been engineered for use in hydrogen environments and rated for sour service by the NACE Standard TM0284 test. All of the steels are fully killed calcium treated continuous cast steels. The X52 has the lowest amount of sulfur

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18 0.0007 wt pct and highest amount of copper 0.35 wt pct. All grades were microalloyed with titanium niobium and/or vanadium to increase strength and control microstructure and transformation products. The thermal mechanical processing and rolling operations were not provided by the steel suppliers. Table 3.1 Chemical Composition of As-Received Plate Steels in wt pct wt pct C Mn Si Ni Cr Mo Ti Nb V Al N S P Cu Ca X52 0.067 1.03 0.24 0.10 0.06 0.03 0.007 0.044 0.007 0.031 0.0082 0.0007 0.010 0.35 0.0027 X60 0.055 1.42 - - 0.08 0.032 Ti+Nb+V – 0.09 - - 0.0028 - - - X70 0.050 1.59 0.30 0.01 0.26 0.09 0.013 0.066 0.005 0.026 0.0081 0.003 0.010 0.01 - 100XF 0.046 1.8 0.21 - - 0.3 0.018 0.070 0.084 0.031 0.009 0.002 0.008 - - “-“ in table represent values that were not reported from the steel supplier upon receiving plate steels 3.2.2 Mechanical Properties Mechanical properties of the plate steels are summarized in Table 3.2. The mechanical properties were obtained from specimens oriented in the transverse direction with respect to rolling direction. The longitudinal and diagonal mechanical properties are not included because prestrain was performed in the transverse direction. Mechanical property data for the 100XF and X70 materials were obtained from Farber 42 and Al Jabr 43 thesis work. X52 and X60 mechanical property data were obtained from the two plate producers SSAB and EVRAZ North America respectively. Table 3.2 Mechanical Properties in the Transverse Direction of As-Received Plate Steels Material 0.2 Offset Yield Strength MPa ksi Tensile Strength MPa ksi YS/UTS pct Elongation 2 in X52 405 58.7 480 69.6 0.86 41.2 X60 515 74.7 585 84.8 0.88 34 X70 465 67.4 595 86.3 0.78 32 100XF 724 105 804 116.6 0.9 18.8 3.2.3 Microstructural Analysis Microstructural analysis on the four plate grades was performed with field-emission scanning electron microscopy FESEM. Faces in the transverse direction TD rolling direction RD and the normal direction ND were polished and etched with 2 pct Nital. The three orientations were evaluated to assess grain shape with respect to the rolling direction the presence of second phase microconstituents and degree of banding. Analyses of polished and etched HIC samples were conducted. Using the energy-dispersive x-ray spectroscopy EDS capabilities of the FESEM the presence and composition of nonmetallic inclusions was identified. Automatic feature analysis AFA was conducted with the aid of Nucor Corporation at their Crawfordsville IN facility. AFA was conducted on two inclusion families conventional melt shop analysis for inclusion control: 1 Ca-Al-S and 2 Mn-Ca-S. Comparison between these two families is a typical melt shop evaluation of the nonmetallic inclusions formed during steelmaking 24. Steels modified for use in hydrogen environments show higher amounts of inclusions in the Ca-Al-S family to increase resistance to HIC. HIC susceptibility is greater if greater amounts of MnS type inclusions are produced during steelmaking. The presence of

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19 MnS is evaluated through the analysis of the Mn-Ca-S inclusion family 24.The purpose of the AFA analysis was to identify the relative composition of the inclusions that formed. The term relative composition is used because the ternary diagrams that were generated represent amounts of each element relative to the other two elements evaluated. Thus when the amount of an element is observed at levels of 90 pct or greater in the ternary diagrams it assumed that the inclusion is comprised of other elements i.e. Fe Mg O etc. 3.2.4 Introduction of Prestrain The experimental plate alloys were pre-strained in tension rather than by rolling or bending in order to avoid strain gradients and residual stresses. A schematic representation of the strip geometry orientation and the prestrain direction of the strips is shown in Figure 3.1. Prestrain was introduced using a 978.6 kN 220 kip tensile machine located at NIST in Boulder Colorado with the help of Mr. D. McColskey. Strips that were 20 mm 0.787 in in width were cut so the specimen longitudinal axis was transverse to the rolling direction. Based on the as-received plates geometry the strip length differed from between each material and was between 450 mm 17.7 in and 700 mm 27.5 in. Prestrain was introduced in the transverse direction with respect to the rolling direction. The presence of Lüder bands in the X52 X60 and X70 materials required specimens from these plates to be subjected to higher amounts of prestrain to ensure uniform strain was introduced throughout the material than specimens from the 100XF plate which did not exhibit such behavior. Strips were first loaded to failure and using crosshead displacement and the grip-to-grip distance of the strip as the gauge length engineering stress-strain curves were generated in order to select displacement/strain values outside the Lüders band region. The prestrain levels for each material and are shown in Table 3.3. Extensometers were not used due to the large grip to grip length of the strips approximately 225 mm 9 in to 305 mm 12 in. To introduce these prestrain levels strips were placed into the load frame and the grip to grip distance was measured as shown in Figure 3.2. This distance was used as the gauge length and the crosshead displacement was then set to a desired distance to obtain the target prestrain level. When the set value of crosshead displacement was reached for each material and prestrain condition the final gauge length was measured to ensure the desired prestrain level had been reached. Prestrain levels for the X60 material were chosen based on recommendations by the steel producer. The X70 material were prestrained to higher amounts than the X60 to evaluate greater amounts of cold work at an increased strength. Prestrain was greatest for X52 material to test if higher amounts of cold work could initiate HIC. The 100XF specimens could not be pre-strained beyond 2 pct because the specimens did not contain a reduced cross-section gauge length and stress concentrations near the gripped sections resulted in non-uniform deformation at relatively small strains. Therefore only one prestrain condition was selected for 100XF.

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20 Figure 3.1 Schematic of strips that were machined in the transverse direction with respect to the rolling direction and then subsequently prestrained to target levels. The prestrain and rolling directions are labeled. Strip length was dependent on the plate material that was machined. The length ranged from 450 mm for the X52 to 700 mm for the 100XF. Figure 3.2 Example of a strip being prestrained using uniaxial tension. The entire length between grips was considered the gauge length. This gauge length was used in correspondence with target prestrain levels to set a predetermined crosshead displacement to achieve the desired prestain amount.

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21 Table 3.3 Selected Prestrain Amounts for the Plate Steels Material Prestrain Amount in pct X52 0 12 18 X60 0 3 5 X70 0 5 7 100XF 0 2 3.2.5 Hardness Traverse and Determination of Microstructural Dependence on Plate Thickness Vickers microhardness measurements were conducted on all as-received plate materials. Cross-sections of the full thickness of the transverse plane were sectioned and polished to 1 µm surface finish. Microhardness measurements were taken every 0.254 mm 0.01 in 500 gmf load and 10 s dwell time from each rolled face edge to the middle thickness of the material. The microhardness traverse data for each as-received material were normalized with respect to plate thickness. For each hardness traverse 50 pct plate thickness represents the mid- thickness of the plate with respect to each rolled face. Edge 1 and edge 2 represent traverses taken from each of the rolled surfaces towards the center of the material. Microhardness measurements were also taken on the prestrained plate materials. Sample size and orientation were kept constant for the hardness traverse procedure. Hardness measurements were taken from the center of the material towards the rolled surface. A minimum of five hardness measurements were performed for each material and prestrain condition additionally care was taken to ensure that no indent was within three diagonals of the previous indent. In order to determine the microstructure as a function of plate thickness micrographs were taken at 20X from the mid-thickness to the rolled face and subsequently pieced together. Micrographs were produced on both non-etched and 2 pct nital etched samples. The non-etched samples illuminate the distribution of nonmetallic inclusions from the rolled surface to the mid-thickness. It should be noted that these micrographs were not used to determine the actual size or composition of the non-metallic inclusions present. The 2 pct nital etched samples were used to identify areas of microstructural differences on a broad scale such as banding and differences in microstructural from the rolled surface to the mid-thickness of the plate. 3.3 HIC Sample Generation and Preparation for HIC Testing The geometry and orientation of the HIC test specimens machined after pre-straining are shown in Figure 3.3. After prestrain was introduced strips were sectioned into shorter pieces approximately 110 mm long. Once these smaller test specimens were created an abrasive saw was used to precisely cut the specimen to the final desired length of 100 mm. To remove the oxide layers on the rolled faces and rust that formed on the cut faces from machining the as-received plates into strips the 100 mm long specimens were placed into the Bridgeport knee mill and face milled. Milling removed 125 hundredths of inch 0.3175 mm from the selected face for each pass until the oxide layer and rust were removed. This procedure was repeated for each of the remaining three faces with respect to normal and longitudinal planes. Milling did not alter the total length of the sample. After milling all faces were

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22 ground and finished with 320 SiC grit paper to maintain a constant surface roughness between test specimens. For specimens that were subjected to electrolytic charging a taper hole was drilled and reamed in the transverse top face of the test specimen to allow for electrical conductivity. After these procedures each test specimen was 100 ± 3 mm 3.937 ± 0.118 in long by 20 ± 3 mm wide 0.7874 ± 0.118 in and the thickness is the full thickness of the plate 34. Three samples of each material and prestrain were fabricated for each hydrogen exposure environment. The test specimen geometry of each specimen created from each material and prestrain condition is found in Tables 3.4 through 3.7. Figure 3.3 Test specimen geometry and orientation used in the NACE Standard and electrolytic charging studies all dimensions are in millimeters. For specimens that were subjected to electrolytic charging a taper hole was drilled and reamed in the transverse top face of the test specimen to allow for electrical conductivity. 3.4 NACE Standard TM0284 H 2 S Method A total of 33 specimens representing the full experimental matrix with 3 replicates for each material and prestrain condition were tested according to NACE standard TM0284 34. The test vessel for the NACE Standard test along with the carousel that the test specimens were placed into is shown in Figure 3.4. Figure 3.4a shows the airtight vessel the test specimens are placed into and the input for hydrogen sulfide gas and the off gas stream. Figure 3.4b shows the carousel that was used to hold the specimens while being submerged in the test solution. The carousel allows for adequate space between the test specimens for equal exposure to the solution and hydrogen sulfide gas. The ratio of volume of test solution to the total exposed surface area of the test specimen was a minimum of 3 mL/cm 2 . Solution A a sodium chloride acetic acid NaCl CH 3 COOH solution saturated with hydrogen sulfide gas at ambient temperature and pressure was used. The specimens were subject to a 96 hour test in this solution after which they were removed and evaluated. Results from the H 2 S method serve a dual purpose in the current study: 1 to validate the experimental procedure for electrolytic charging to produce HIC and 2 to

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23 investigate the effects of prestrain on HIC susceptibility. This test was conducted with the help of EVRAZ North America at their Research and Development facilities located in Regina Saskatchewan. Table 3.4 Test Specimen Geometry for the X52 Specimens used in each Hydrogen Exposure Method NACE Standard test specimen geometry Material Prestrain Specimen No. Width mm Thickness mm Length mm X52 0 1 19.17 18.38 101.48 2 19.17 18.31 101.78 3 19.23 18.4 101.19 X52 12 1 18.13 17.26 100.72 2 18.45 17.22 101.02 3 18.34 17.23 101.04 X52 18 1 17.67 16.84 100.98 2 17.52 16.82 100.85 3 17.48 16.93 100.17 EC test specimen geometry Material Prestrain Specimen No. Width mm Thickness mm Length mm X52 0 1 19.23 18.31 99.89 2 19.37 18.21 101.21 3 19.41 25.58 102.99 X52 12 1 17.93 17.78 102.77 2 17.91 17.42 102.92 3 17.81 17.32 98.54 X52 18 1 17.93 16.84 102.33 2 17.81 16.99 100.82 3 17.78 16.69 102.76

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24 Table 3.5 Test Specimen Geometry for the X60 Specimens used in each Hydrogen Exposure Method NACE Standard test specimen geometry Material Prestrain Specimen No. Width mm Thickness mm Length mm X60 0 1 19.4 8.24 98.3 2 19.82 8.49 98.77 3 19.28 8.36 98.51 X60 3 1 19.56 8.66 99.74 2 19.88 8.68 99.55 3 19.68 8.4 98.52 X60 5 1 19.48 8.92 99.38 2 19.7 8.86 98.75 3 19.73 8.9 99.42 EC test specimen geometry Material Prestrain Specimen No. Width mm Thickness mm Length mm X60 0 1 20.04 8.89 98.97 2 19.66 8.92 102.21 3 19.69 9.02 100.71 X60 3 1 20.12 8.69 98.14 2 19.81 8.69 101.26 3 19.71 8.69 98.62 X60 5 1 19.91 8.53 100.35 2 20.12 8.33 100.63 3 20.08 8.28 100.87

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25 Table 3.6 Test Specimen Geometry for the X70 Specimens used in each Hydrogen Exposure Method NACE Standard test specimen geometry Material Prestrain Specimen No. Width mm Thickness mm Length mm X70 0 1 19.48 12.1 100.01 2 18.96 12.13 99.91 3 19.01 12.06 99.35 X70 5 1 19.05 11.83 100.01 2 18.97 11.84 99.74 3 19.07 11.91 99.78 X70 7 1 18.8 11.89 99.75 2 18.97 11.87 100.69 3 18.57 11.8 99.7 EC test specimen geometry Material Prestrain Specimen No. Width mm Thickness mm Length mm X70 0 1 19.58 12.12 100.73 2 19.42 12.17 100.49 3 19.61 12.07 100.44 X70 5 1 19.1 11.84 100.22 2 19.08 11.58 100.61 3 19.23 11.89 99.84 X70 7 1 19.08 11.74 99.62 2 19.03 11.68 102.73 3 19.03 11.71 98.63

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26 Table 3.7 Test Specimen Geometry for the 100XF Specimens used in each Hydrogen Exposure Method NACE Standard test specimen geometry Material Prestrain Specimen No. Width mm Thickness mm Length mm 100XF 0 1 19.26 11.77 100.47 2 19.39 12.11 100.16 3 18.78 11.85 100.15 100XF 2 1 18.72 12.08 98.11 2 18.77 12.1 98.65 3 18.74 12.09 98.27 EC test specimen geometry Material Prestrain Specimen No. Width mm Thickness mm Length mm 100XF 0 1 18.69 12.27 101.63 2 19.38 12.37 100.83 3 19.28 12.02 98.11 100XF 2 1 19.43 12.34 99.93 2 19.33 12.16 100.88 3 19.48 12.34 100.58 a b Figure 3.4 a NACE Standard TM0284 test vessel used to evaluate materials suscpetibilty to HIC b carousel used to hold test specimens allowing for adequate space for gas purge during testing. 3.5 Electrolytic Charging Methodology EC The following section highlights the research and development of the EC methodology. Design issues encountered from previous attempts at EC charging for HIC resistance are presented. Then solutions to address

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27 these design issues are discussed as well as the fabrication of the new EC apparatus used in the present study. A characterization study of the new EC cell was undertaken to optimize the experiment for HIC assessment. Based on the characterization study an experimental procedure was developed and implemented to investigate HIC using EC. 3.5.1 Initial Design and Fabrication of New Charging Apparatus Cathodic charging to introduce hydrogen to the test specimens has been selected as a substitute for NACE standard test TM0284 34. Figure 3.5 shows the previous generation cell used for electrolytic charging. The cell is a simple beaker referred to as a monocell that holds the electrolyte and both electrodes steel cathode and graphite anode. In previous attempts to validate this electrolytic charging setup for long tests e.g. 24 hours or greater it was observed that the current within the system tends to vary when the system is voltage controlled or voltage varies if the applied current is controlled. The slight variation in the electrical response of the system is important to monitor in order to assess the reliability of the charging. Due to the long term nature of the charging Rosner 7 and Angus 44 discovered issues related to the stability and repeatability of electrolytic charging experiments. The issues Rosner 7 and Angus 44 summarized and the solutions to them are summarized in Table 3.8. Table 3.8 Design Issues and Solutions taken for the Electrolytic Charging Methodology Design Issue Design Solution During electrolytic charging the anode material graphite underwent mass loss and changes in geometry. The mass loss of the graphite would then pollute the electrolyte changing the nature of the charging conditions within the cell. A titanium mesh that was coated in ruthenium oxide provided by Dr. Martins was selected as the anode. This material does not undergo the weight loss and changes in specimen geometry during testing as observed for the graphite. Alligator clips have been used as the main method to achieve an electrical connection to the cathode which is the steel specimen. At long test durations i.e. greater than 18 hours the alligator clips would fail and the electrical connection to the steel specimen would be lost limiting the possible test duration. A new taper pin plug in friction fitted electrode connection was devised. The taper pin fit to a predrilled hole into the steel specimen to create the electrical connection. To protect this connection further the area of the taper pin exposed to the aqueous solution was coated in epoxy. The initial trials for electrochemical hydrogen charging of test specimens in aqueous solutions used a single cell design. Both the anode and cathode were in the same compartment. This lead to cross-contamination of ion species at each anode/cathode interface. This cross- contamination causes unwanted reactions to occur at those interfaces lowering the overall effectiveness of electrolytic charging. It was also observed that corrosion and mass loss of the steel specimen could occur in this cell due to increased concentration of ions within the system around the steel specimen regardless of the cathodic overpotential that the steel specimen was experiencing. To minimize cross-contamination within the system a newly designed dual cell was fabricated. The dual cell has separate anodic and cathodic compartments. Ion separation was achieved through a salt bridge between the compartments. This salt bridge allows for anodic species in the cathodic compartment to flow to the anodic compartment and vice versa for cathodic species. Purge gas systems was fabricated for each compartment to aid in the flow of ionic species from one compartment to another and minimize potential ionic concentration gradients. Oxide scale from the steel specimen flaked off during testing and polluted the electrolyte. This pollution altered the nature of the electrolyte and changed to electrical response within the system. Each steel specimen face was milled and subsequently ground to a final surface roughness of 320 grit to remove oxide scale.

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28 Figure 3.5 Prior generation cell used for electrolytic charging by B. Rosner 7 and G. Angus 44. The initial cell in Figure 3.5 was a very simple monocell design. In order to achieve ion separation the anodic and cathodic compartments were separated to create a dual cell design schematically shown in Figure 3.7. To allow for ion species to travel from one compartment to another the two compartments are connected to each other via a side branch with a salt bridge. An O-ring seal is used at the connection which is fastened together with flanges. Each compartment is fitted with a Teflon cover that incorporates ports with Teflon compression-fittings for positioning of each electrode assembly as well as for gas effluent-discharge. Fitted aluminum flanges provide a means to attach the Teflon compression-fittings to each compartment. A side-connection on the bottom of the main body of each compartment allows inert-gas purging of the anodic and cathodic compartments respectively. Altering the initial test apparatus to the proposed dual cell design aids in stabilizing EC experiments at prolonged test durations. Using the design solutions in Table 3.8 a newly designed test cell for electrolytic charging was fabricated. The anode material was changed from the graphite anode to the titanium mesh coated in ruthenium oxide and this assembly is referred to as a dimensionally stable anode DSA. The anode counter-electrode assembly was fabricated at Colorado School of Mines and consisted of a titanium rod welded to a titanium collar that is welded to the DSA. The composite material comprising the DSA is a diamond-pattern expanded titanium-metal coated with ruthenium oxide semi-conductor. The electrical connector to the working-electrode test-specimen consisted of a hardened taper-pin welded to a 316SS rod. Fabrication of the cathodic electrical connector is shown in Figure 3.6a. A taper hole to accommodate this connector was drilled and reamed in the transverse top face of the test specimen shown in Figure 3.6b. Once

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29 inserted and friction locked into the test specimen the lower region of the connector was electrically insulated with a thin coating of epoxy resin. a b Figure 3.6 From left to right the images display the salient features for a the construction of the cathode electrode connection to be used in correspondence with b the steel specimens. The cathode electrode connection in a employs a taper pin fastened to a 316 SS rod. A corresponding taper hole is reamed into the top transverse face of the steel specimen. The cathode electrode connection is then inserted into the hole and frictionally locked with the test specimen. To protect this electrical connection the region around the connection is covered in epoxy. Figure 3.7 Schematic diagram for dual cell used for electrolytic charging experiments. The fabricated EC cell is shown in Figure 3.8. The separate anodic and cathodic compartments were custom made Pyrex glass vessels fabricated by Allen Scientific Glass Inc. of Boulder Colorado. The fitted Teflon

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30 compression fittings were fabricated by Dr. Martins. The two compartments when assembled are held by a custom- designed aluminum stand fabricated by Stephen Tate. The gas purge ports include a Teflon up-leg connection to allow for inert-gas purging argon gas fabricated by Dr. Martins. The flow rate of the gas is controlled by a flow meter found in Figure 3.8. A BK Precision 1735A 30V/3A DC power supply with high sensitivity voltage ≤ 0.02 ± 3 mV current ≤ 0.2 ± 3 mA was used to supply the current the Potentiostat found in Figure 3.8. 3.5.2 Characterization Study and Experimental Procedure for EC Electrolytic charging EC was accomplished through electrochemical polarization using an electrochemical cell designed to optimize the reproducibility of the HIC results and the stability of the components during long charging times. Current was controlled using a Potentiostat while voltage varied based on the electrochemical nature of reactions occurring between the two electrodes in the test setup. A characterization study on the newly fabricated test cell Figure 3.8 for EC was undertaken. The 100XF as-received plate material was chosen for this study based on the assumption that it would show the highest susceptibility to HIC and facilitate identifying the parameters that would alter the reliability of EC. The parameters that were altered during this study were:  The applied current voltage density  Test duration Applied current density and test duration were chosen as variables to alter within the system based on a study by Perez Escobar et al. 41. The results Perez Escobar et al. obtained showed increased levels of hydrogen damage were caused from increases in both applied current density and test duration. Changes in electrolyte chemistry were not explored in this study because these effects have been explored in previous theses 6 and 7. Other variables such as flow rate of gas purge cathodic surface exposure to gas purge bubble stream temperature of electrolyte and surface finish of steel specimen were held constant throughout the characterization study. Tests were conducted at room temperature 20 – 22 °C and ambient barometric-pressure 0.80 atm. The surface finish of the steel specimen was ground to 320 grit as described in section 3.3 to avoid contamination of electrolyte during testing. The results from the characterization study are presented in section 4.3. The following experimental procedure for EC was established based on this study for the remaining investigation of HIC in the material and prestrain conditions outlined in Section 3.2:  Current density of 15 mA/cm 2  Test duration set at 24 hours  The total volume of the electrolyte maintained at 750 mL  The argon gas purge set to 25 cm 3 /min  The electrolyte is 1 normal H 2 SO 4 solution with an addition of 20 mg/L of As 2 O 3  Surface finish of steel specimen ground to 320 grit  No manual alteration to increase temperature during testing tests were conducted at room temperature  No alteration to pressure within the test apparatus was undertaken tests were conducted at ambient barometric pressure

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31 Figure 3.8 Electrolytic hydrogen charging EC test apparatus with components labeled. Electrolytic charging EC was accomplished through electrochemical polarization using this test apparatus. Tests apparatus was designed to optimize the reproducibility of the HIC results and the stability of the components during long charging times

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32 3.6 Assessment of Hydrogen Damage Upon completion of each hydrogen charging method H 2 S or EC the test specimens were taken out of the solution and cleaned and the electrical connection was subsequently removed from the charged specimens. Each specimen was then sectioned as shown in Figure 3.9a as per the NACE Standard Test only faces 2 4 and 6 were evaluated for specimens that were charged by the H 2 S method whereas all 6 locations were evaluated for specimens charged by the EC method. Light optical microscopy was used to determine the NACE crack ratios on the polished sections as shown schematically by Figure 3.9b. The Crack Sensitivity Ratio CSR Crack Length Ratio CLR and Crack Thickness Ratio CTR were calculated with the following equations: where a is length of single crack b is the thickness of a single crack including crack branching and W and T are the width and thickness of the cross-section respectively. Cracks separated by less than 0.5 mm were considered as a single crack. All cracks observable at magnifications as high as 100X are included in the three calculations. The average of each ratio was calculated for all of the examined faces of each test specimen 34. a b Figure 3.9 Test specimen geometry and orientation used in the NACE Standard and electrolytic charging studies all dimensions are in millimeters a Faces 2 4 and 6 are used in the NACE Standard evaluation while Faces 1 - 6 are used for the electrolytic charging method. Face 1 is the furthest from the electrode link. b Schematic showing how cracks are measured on the evaluated faces 34. ∑ ⁄ 3.1 ∑ ⁄ 3.2 ∑ ⁄ 3.3

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33 The hydrogen charged specimens from both the EC method and NACE standard test were evaluated according to the procedures outlined in NACE standard test TM0284 34. The specimens from the NACE standard test conducted at EVRAZ were evaluated by a technician and the results were provided in an Excel spreadsheet report. The method used to measure cracks on the EC cross-sections started by taking an image of the full cross- section of the examined face using the stereoscope the image was imported into ImageJ an image processing program. . The scale of the image was set in ImageJ. A screen shot of this process is shown in Figure 3.10. Figure 3.10 shows the interface from the program along with the measurement of the plate thickness in order to set the scale. Once the plate thickness is measured and the corresponding value in pixels is found the “set scale” feature in the program the amount of pixels can be set to the known plate thickness in mm. For the image shown in Figure 3.10 the value in pixels ‘502’ corresponds to 12.42 mm. ImageJ also displays the resolution of the image in reference to pixels which was 40.419 pixels per mm in Figure 3.10. Taking the reciprocal of this value to get mm per pixel it is assumed that calculations of crack lengths and thickness from the image are accurate out to 3 hundredths of an mm 30 µm. Figure 3.10 ImageJ interface that allows the user to import selected images into the software for analysis. Once imported the plate thickness is measured in pixels shown by the white line. Knowing this distance in pixels the “set scale” feature in the software allows the user to set the distance in pixels to a known distance. For this image and cross-section 502 pixels equaled 12.42 mm. The “set scale” feature also displays the resolution of the image relative to the scale. This resolution was used to determine the accuracy of the measurements. Cracks were identified and measured using the steps shown in Figure 3.12. A crack was distinguished from a scratch by the jagged features of the cracks compared to the straight line morphology of the scratches. According to the NACE standard “cracks separated by less than 0.5 mm are considered as a single crack 34.” An example of

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34 the measured separation in cracks is shown in Figure 3.12a. Distances between the cracks were measured as shown in Figure 3.12a. Because the offset of cracks identified in regions No. 1 2 and 6 in Figure 3.12a were less than 0.5 mm Figure 3.12b the two cracks in each region are considered a single crack. Cracks identified as No. 3 – 5 Figure 3.12b have an offset distance greater than 0.5 mm so crack was identified as a single crack. Using the “zoom” feature in ImageJ each crack is enhanced to measurement length a and thickness b of the crack Figure 3.12c. White bars represent the measurements of the 8 identified cracks shown in Figure 3.12c. The corresponding values of length and thickness for each crack are shown as an example of representative measurements in Table 3.9. Using Equations 3.1 – 3.3 the crack ratios were calculated Table 3.9. Table 3.9 Measured values from Manual Image J Measurements of length ‘a’ and thickness ‘b’ of Cracks in Figure 3.12. The Calculated Crack Ratio Values are also shown. Material Prestrain condition 100XF 0 pct W 24.77 mm T 12.47 mm Crack a mm b mm CLR CTR CSR 1 1.93 0.29 fractional fractional fractional 2 11.93 1.06 1.1114 0.2013 0.0524 3 2.23 0.07 4 1.65 0.07 5 2.05 0.07 6 4.11 0.21 7 2.59 0.61 8 1.04 0.13 The uncertainties for the measurements used in equations 3.1 – 3.3 are shown in Table 3.10. The uncertainty for variables a and b are in reference to the image resolution imported into Image J used to measure each individual crack. Values of uncertainty for the W and T variables are from the calipers that were used to measure the cross-section of the selected face. The total uncertainty for equations 3.1 – 3.3 are calculated by the root sum of squares equation where w x is the uncertainty in the device used to obtain the measurement and δCTR CLR CSR/δx is the partial derivative of the function with respect to the variable associated with the measuring device. Using Equations 3.4 and 3.5 values of total uncertainty were calculated for each plate material. The uncertainty values for each material and critical crack ratio are shown in Table 3.10. All uncertainty values for critical crack ratios are accurate out to approximately 0.0005 of a fractional value except the CTR value for the X60. Because the X60 is the thinnest plate material out of the four the uncertainty is 0.001 of a fractional value. ∑ ⁄ 3.4 ∑ ⁄ 3.5

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35 Table 3.10 Uncertainty from the Method used to Measure Variables for use in Crack Ratio Calculation Method Variable Uncertainty mm Caliper W and T 0.005 Image J a and b 0.03 3.7 LECO® Hydrogen Analysis The LECO® RH-404 hydrogen analyzer induction melts a test specimen and reports the total hydrogen content at the time of measurement in units of ppm mass fraction. For accurate measurements of hydrogen content the weight of the specimens that are placed into the hydrogen analyzer should be 1.0 ± 0.5 g. Total hydrogen levels in the current study were measured using sub-size test specimens that were placed into the LECO® RH-404. As- received plate materials were evaluated for hydrogen content to serve as a baseline of hydrogen present in the material before charging. Total hydrogen levels were also measured for each material and pre-strain condition. 3.7.1 LECO® Hydrogen Analysis Sample Generation A means to charge multiple sub-sized specimens was generated because the hydrogen analyzer is only able to analyze small specimens. The specimen geometry is set to: plate thickness which ranges from 0.95 to 1.9 cm x 0.2 x 0.7 cm 0.375 to 0.75 in x 0.079 x 0.276 in. Figure 3.11 shows the fabrication of a “daisy chain” that allowed for multiple sub-sized specimens to be charged for hydrogen evaluation. For each material and prestrain condition a total of 5 replicas were created. Samples of each material and prestrain condition were cut to the geometry mentioned above and using copper wire wound so that they were securely in contact with the wire to ensure electrical conductivity. Figure 3.11 “Daisy Chain” assembly components. From left to right the images display the salient features the five sub-sized specimens and copper wire of the component assembly leading to the assembly of the “Daisy Chain” to allow for simultaneous charging of multiple specimens for hydrogen content measurements.

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36 a b c Figure 3.12 Example cracked face showing how images were used to a identify cracks that are offset from one another. b Example cracked face showing the measurement offset to determine if multiple cracks should be considered a single cracks separated by less than 0.5 mm were considered as a single crack 34. c The final measurement of the single cracks. Each crack in b was enhanced using the “zoom” feature to produce the images in c to aid in the precise measurement of the single crack. The white lines in c represent the length a and thickness b measurements used to calculate critical crack ratios.

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37 3.7.2 LECO® Hydrogen Analysis Experimental Procedure With “daisy-chains” fabricated for each material and prestrain condition Figure 3.11 the total exposed surface area of the samples and copper wire was determined to apply the correct current at the applied current density of 15 mA/cm 2 . The “daisy-chain” was placed into the EC apparatus Figure 3.8 and charged with hydrogen for 24 hours following the sample procedure outline for full size test specimens that were evaluated for HIC. After charging the “daisy chains” were removed cleaned with water and placed into a liquid bath of nitrogen within 30 seconds in order to suppress the diffusion of hydrogen out of the samples. LECO® steel standards with known amounts of hydrogen were then used to calibrate the H analyzer before measuring the steel specimen hydrogen content. The resolution of the LECO® RH-404 is 0.05 ppm H. After calibrating the LECO® H analyzer individual samples were taken from the bath of liquid nitrogen removed from the “daisy chain” rinsed dried and placed into the LECO® RH-404 for hydrogen analysis. For each sub-sized test specimen the time spent at room temperature 22 °C before the LECO® RH-404 induction melted and evaluated the hydrogen content was on average 90 seconds. Diffusible hydrogen concentration for the LECO® method were calculated by where H Total is the hydrogen concentration determined by the LECO® immediately after EC and H Trapped is the hydrogen concentration evaluated after degassing determined by the LECO® analyzer. Degassing was achieved by placing full size test specimens in a liquid water bath at an elevated temperature of 50 °C for 72 hours as described in Section 3.8. Samples were sectioned from these degassed full size test specimens to create the sub-sized specimens needed for LECO® analysis. A total of 5 specimens were generated from each material and prestrain condition. The 5 values from each hydrogen concentration value H Total and H Trapped were then used in combination with Equation 3.6 to produce LECO Diffusible values. 3.8 Diffusible Hydrogen Content Determined by Mercury Displacement The American Welding Society AWS created a standard for diffusible hydrogen measurements based on mercury displacement. This section presents the experimental procedure that was modified from this standard AWS A4.3-86 45 and implemented to measure diffusible hydrogen contents after exposure to the EC method. 3.8.1 Mercury Displacement Sample Generation Multiple specimens for diffusible hydrogen analysis were created from a single sample charged with hydrogen by EC. A visual representation of the fabrication of these sectioned EC test specimens is shown in Figure 3.13. Full size test samples were sectioned so that after EC trials 4 equal sized specimens test samples labeled 1 2 3 and 4 in Figure 3.13 could be generated without further processing steps. After the taper hole was drilled into the top transverse section the specimen was placed into the abrasive saw where it was partially sectioned at the 25 50 and 75 mm locations Figure 3.13a. The black area in Figure 3.13a represents the material left after sectioning visually shown in Figure 3.13b designated as the areas of electrical connection by the arrows. The 3.6

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38 overall dimensions of the cut were not recorded but care was taken to leave enough material at each section location to ensure electrical conductivity but remove enough so sections could be separated from one another after EC trails. The sectioning of these areas produced 6 added surfaces that were exposed to solution and EC. The surface area of these surfaces was factored into the current calculation calculated to produce an applied current density of 15 mA/cm 2 . It was assumed that the material left from the sectioning process was minimal such that the total width and plate thickness dimensions were used to calculate the additional exposed surface area. The 4 test specimens generated from this method allowed for an average diffusible hydrogen concentration to be calculated as well as the distribution of diffusible hydrogen along the specimen i.e. from test sample 1 to sample 4. a b Figure 3.13 a Schematic representation with specimen geometry and the locations where test specimens were sectioned. The black areas represent material that was left after the sectioning process was complete. b Visual representation of samples fabricated for diffusible hydrogen analysis using EC. The sectioned areas allow for easy detachment after EC has been conducted. 3.8.2 Mercury Displacement Experimental Procedure The procedure outlined by the AWS 46 covers correct handling storage and placement of specimens into the mercury filled eudiometers for measurement as well as incubation periods and the final measurement of diffusible hydrogen. The following section highlights important details from the AWS standard pertinent to setting up the procedure for the current study. The specimen of interest is required to be placed into a low-temperature liquid bath -60 °C -76 °F within 60 seconds after removal from the hydrogen environment. After test specimens have been in this liquid bath for two minutes they may be removed for further handling i.e. cleaning detaching specimens from one another and/or being placed into the mercury displacement apparatus 45. It is stated that when removed from the low-temperature bath a specimen has to be placed back into the bath after one minute exposure to ambient conditions and only after

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39 two minutes in the bath may it be removed again 45. This procedure is used each time a specimen is removed from the liquid nitrogen bath to ensure that hydrogen remains trapped before further analysis. Diffusible hydrogen analysis is performed with a mercury-filled eudiometer. Figure 3.14 shows schematically the dimensions of the eudiometer Figure 3.14a as well as the evolution of hydrogen from a sample placed into the mercury-filled eudiometer Figure 3.14b. The test apparatus requires loading a test specimen into the incubation chamber located at the lower end of the eudiometer allowing for hydrogen to diffuse out of the sample and recombine into molecular hydrogen. The process of removing the test specimen from the low- temperature bath and placing into the incubation chamber is advised to be within 150 seconds from start to finish to ensure reliable results. During the degassing of the sample in the eudiometer hydrogen displaces the mercury away from the top valve the Teflon stopcock towards the mercury bath at the bottom. Because there is a physical height difference of mercury in the eudiometer and the lower mercury bath there is a pressure difference for the evolved hydrogen compared to the ambient pressure on the mercury bath. When calculating the volume of hydrogen this pressure difference is incorporated. The time required before evaluation is dependent on the liquid bath temperature. After the incubation period the volume of hydrogen present and height of displacement are measured to calculate diffusible hydrogen value at STP. The volume of hydrogen gas is calculated as where T is the temperature °C of the gas column at the time of measurement room temperature P is the barometric pressure mm Hg at the time of measurement V is the measured volume mL taken from the eudiometer H is the head height of mercury mm at time of measurement as shown in Figure 3.14b and VH is the volume of hydrogen gas at STP in milliliters 45. The same conditions used to generate HIC by EC were employed to test for diffusible hydrogen in the sectioned specimens for each alloy and prestrain conditions. After EC was completed the steel specimen was removed from the EC test apparatus rinsed with water and placed into a low-temperature liquid nitrogen bath within 60 seconds. The specimens were transported to the displacement test apparatus. The mercury test setup is shown in its entirety in Figure 3.15 shown are the four mercury-filled eudiometer tubes the temperature setting and thermometer for the water bath and the ambient pressure gauge. When placing the samples into the eudiometers the following procedure was implemented:  Remove full-size test specimen from liquid nitrogen bath  Using pillars break sample of interest from the full-size test specimen placing remaining pieces back into the bath of liquid nitrogen  Clean in liquid water bath to remove layer of ice created by exposure to the atmosphere  Dry sample so no moisture is present on the surface failing to do so will give inflated readings of VH  Place sample into the bottom end of the eudiometer this assembly is then inverted and placed into the mercury 3.7

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40  The mercury is drawn up to the top of the eudiometer with a vacuum after which the top valve the Teflon stopcock is closed to allow for hydrogen to be captured and displace the mercury Figure 3.14 a Dimensions of the eudiometer used in the mercury displacement method to determine diffusible hydrogen amount b schematic representation of the use of a mercury-filled eudiometer to capture and measure the amount of diffusible hydrogen in the sample that is placed into the assembly. Adapted from 45.

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41 Figure 3.15 Test setup for the mecury displacement method with the compontents labeled. For the current study the temperature of the liquid bath was set to 50 °C which corresponds to a 72 hour incubation period. After the incubation period is completed the height volume and other required variables at time of measurement are recorded such that VH Equation 3.7 can be calculated for each specimen. In order to compare the measured hydrogen contents with those obtained through other hydrogen analysis techniques the VH value was converted to parts per million. The parts per million of diffusible hydrogen for each sample was calculated as where VH is the calculated value from Equation 3.7 m steel is the mass of the individual steel specimen is the density of molecular hydrogen gas at STP and 10 -6 is the conversion factor to parts per million. Uncertainty for the mercury displacement method was calculated by the root sum of squares equation where wH is the uncertainty for the height measurement taken by a tape ruler with a resolution of 1 cm wM is the uncertainty from the analytical balance used to measure the mass of the individual samples wP is the uncertainty in the ambient pressure gauge wT is the uncertainty in the thermometer used to determine room temperature wV is the uncertainty of the volume taken from the eudiometer and are the partial derivatives of Equation 3.8 with respect to each variable that was 3.8 wppm 3.9

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42 measured. Uncertainty values for Equations 3.7 and 3.8 are shown in Table 3.11. Using a range of values of diffusible hydrogen from results presented in section 4.5 total uncertainty calculations were made for each material using Equation 3.9 and values of uncertainty in Table 3.11. The results for total uncertainty are shown in Table 3.12. Lower amounts of diffusible hydrogen in the X52 produced lower amounts of uncertainty 0.02 ppm and at the highest value of diffusible hydrogen measured in the 100XF the uncertainty increased to 0.05 ppm. Table 3.11 Uncertainty Values for the Instruments Used to Calculate Diffusible Hydrogen Method Variable Uncertainty Eudiometer V 0.01 mL Pressure gauge P 0.01 mm Hg Tape ruler H 5 mm Analytical balance m steel 0.01 grams Standard thermometer T 0.01 °C Table 3.12 Total Uncertainty Calculated using Equation 3.9 for the Mercury Displacement Method Material Method Variable Diffusible hydrogen ppm Total uncertainty ppm X52 Mercury Displacement ppm 1.56 0.02 X60 Mercury Displacement ppm 3.99 0.04 X70 Mercury Displacement ppm 3.97 0.04 100XF Mercury Displacement ppm 6.42 0.05

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43 CHAPTER 4: Results and Discussion This chapter will present and discuss the results obtained from the different experimental methods discussed in the previous chapter. 4.1 Microstructure and Nonmetallic Inclusions The microstructure of the as-received materials was evaluated using standard metallographic techniques. Faces in the transverse direction TD rolling direction RD and the normal direction ND were polished and etched 2 pct Nital. The X52 plate material was produced by controlled rolling and then cooled using on-line accelerated cooling. FESEM micrographs of X52 are shown in Figure 4.1. Micrographs were taken on the three different planes to evaluate the uniformity of the microstructure. The micrographs exhibit a mixed ferrite/pearlite structure. The ferrite in the microstructure is polygonal or equiaxed. Figure 4.2 shows higher resolution image of the secondary phase that formed in the X52. The pearlite Figure 4.2 in the microstructure is not fully developed and does not resemble the conventional lamellar morphology caused by the low carbon levels hypoeutectoid and the cooling operations performed on the X52. It is described as degenerated pearlite which forms due to insufficient carbon diffusion when the eutectoid reaction proceeds at a low temperature 46. The degree of elongation with respect to rolling direction is minimal i.e. minimal pancaking of grains and the degree of banding is low. The X60 and X70 material were both produced from continuous cast slabs that were subsequently hot- rolled into plate. The micrographs of the three planes with respect to rolling direction are shown in Figures 4.3 and 4.4 for the X60 and X70 respectively. The difference in microstructures is minimal as shown in Figures 4.3 and 4.4. The microstructure is dominated by acicular ferrite with some quasi-polygonal ferrite. The presence of second phase microconstituents martensitic islands appears to be slightly greater in the X70 as shown by the circled regions in each micrograph. Rolling operations and cooling practices for these plate materials resulted in a small amount of elongation of grains with respect to the rolling direction. The three micrographs from each plane with respect to rolling direction for the 100XF steel are shown in Figure 4.5. The plate was produced by controlled rolling followed by accelerated cooling which is expected to result in an acicular ferrite microstructure 42. The micrographs show elongated grains as well as some that are more equiaxed in shape. The microstructure can be described as acicular ferrite with minimal quasi-polygonal ferrite. The 100XF exhibits relatively higher amounts visual comparison of second phase microconstituents indicated by the black circled regions in the micrographs compared to the other alloys. The plates selected in the current study had calcium additions to aid in controlling the types composition and morphologies of nonmetallic inclusions present. As mentioned in the materials selection section the X52 and X60 materials were produced for use in hydrogen environments. A general guide for steels used in sour service is to have ratios of Ca:S at levels greater than 2 24. It is expected that the addition of calcium resulted in inclusions that are globular instead of elongated and the amount of MnS type inclusions is minimized. During steelmaking the

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44 aluminum and calcium react with the oxygen and sulfur before manganese or other elements in the melt therefore minimizing the formation of MnS type inclusions 24. a b c Figure 4.1 Secondary electron micrographs taken with the FESEM of X52 plate steel. Images from three different planes a transverse b longitudinal and c normal plane are shown. 2 pct Nital etch. Figure 4.2 Presence of degenerated pearlite in the X52 plate material. Image taken with the FESEM. 2 pct nital etch.

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45 a b c Figure 4.3 Secondary electron micrographs taken with the FESEM of X60 plate steel. Images from three different a transverse b longitudinal and c normal plane planes are shown. 2 pct nital etch. Black circles on each micrograph indicate the presence of a secondary microconstituent. a b c Figure 4.4 Secondary electron micrographs taken with the FESEM of X70 plate steel. Images from three different planes a transverse b longitudinal and c normal plane are shown. 2 pct nital etch. Black circles on each micrograph indicate the presence of a secondary microconstituent. a b c Figure 4.5 Secondary electron micrographs taken with the FESEM of 100XF plate steel. Images from three different planes a transverse b longitudinal and c normal plane are shown. 2 pct nital etch. Black circles on each micrograph indicate the presence of a secondary microconstituent.

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46 Nonmetallic inclusions in each steel are shown in the images of Figure 4.6. The micrographs and their corresponding EDS maps represent the common shape size and compositions of the inclusions that were found. Because an extensive statistical inclusion study was not conducted the possibility of other types sizes and morphologies cannot be ruled out. Metallographic observations of the alloys showed that the inclusions were 6 ± 1 µm 2 on average and were largely globular in their morphology Figure 4.6. Several inclusion types were found including: Al-Ca-O Figure 4.6a Al-O Figure 4.6b Al-Mg-O-Ca-S Figure 4.6c and Al-Mg-O Figure 4.6d. The composition of inclusions presented here were present within each steel evaluated. a b c d Figure 4.6 Nonmetallic inclusions observed in the four plate steels. Elements present in each image were confirmed by EDS mapping of the image shown. a Al-Ca-O X52 b Al-O X60 c Al-Mg-O- Ca-S X70 d Al-Mg-O 100XF see pdf version for color. The ternary diagrams produced for X52 from two inclusion families 1 Al Ca S and 2 Ca Mn S are shown in Figure 4.7a and b respectively. The values of composition determined by the tie lines in the ternary diagram are relative to the other elements present not necessarily the actual composition of inclusion evaluated. The grouping of compositions in the two ternary plots is an example of the qualitative behavior modified steel for use in sour service would display 24. Region 1 in Figure 4.7a shows inclusions in the family of Al-Ca-S on average have relative levels of Ca at 40 to 80 wt pct S at levels of 50 wt pct or less and Al at a maximum of 30 wt pct . Lower amounts of sulfur in the inclusion compositions are expected because of the ultra-low level of sulfur 0.0007 wt pct that was achieved during plate production. Figure 4.7b shows that the amount of MnS inclusions formed in the X52 is minimal. Most of the inclusions in the Ca-Mn-S family appear to be CaS. This is displayed by the dashed region 2 in Figure 4.7b: Ca levels range from 30 to 85 wt pct S from 15 to 70 wt pct and Mn from 0 to 20 wt pct. Figure 4.8 shows the ternary diagrams for the X60 inclusions in the Ca-Al-S and Ca-Mn-S families. Dashed regions in Figure 4.8a and b represent the common range of inclusion compositions within each ternary

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47 diagram. As shown in the dashed regions 1 and 2 in Figure 4.8a and b the inclusions in the X60 plate material have a higher tendency compared to the X52 Figure 4.7b to be rich in sulfur at least 30 wt pct or greater. The X60 also exhibits higher propensity to form aluminum rich inclusions shown in the dashed region 1 of Figure 4.8a. There is also an increase in the presence of manganese in the inclusions in comparison to the X52 with a large number of values ranging from 0 – 50 wt pct shown in the dashed region 2 of Figure 4.8b. The level of manganese added to the X60 is 1.42 wt pct compared to 1.03 wt pct for the X52. The higher manganese content present in the X60 increases the likelihood for more MnS type inclusions to form with unreacted sulfur in the melt causing the shift in Ca-Mn-S ternary plot towards the manganese sulfur tie line Figure 4.8b. It was expected based on literature findings 3 10 12 15 22 and 23 that cracks generated in this material will originate around MnS inclusions. a b Figure 4.7 Ternary diagrams in wt pct of two inclusions families a Al-Ca-S and b Ca-Mn-S present in the X52 plate steel. Ternary diagrams were generated by automatic feature analysis at the Nucor Indiana facility. Dashed regions 1and 2 in a and b respectively show the grouping of the relative compositional distribution of inclusions evaluated by AFA in reference to the three elements found in each ternary diagram. The analysis of the two inclusional families 1Al-Ca-S and 2 Ca-Mn-S for the X70 are shown in Figure 4.9. Dashed regions 1 and 2 in Figure 4.9a and b represent the common range of inclusion compositions within each ternary diagram produced from AFA. The bulk of inclusions present in the X70 material are greater in calcium Figure 4.9a and b than the levels found in the X60. Like the X60 there is more scatter in the compositions of the inclusions when compared to the X52 material as indicated by the qualitative assessment of the areas encompassed by the two dashed regions. The relative amount of inclusions with the composition rich in manganese is comparable to the X60 Figure 4.8b but greater than the X52 Figure 4.7b. This is expected from the amount of manganese present 1.59 wt pct. The X70 compositions indicate the presence of MnS regions 3 in Figure 4.9b. The formation of MnS is related to ladle metallurgy practices 24. It was expected that cracks generated through hydrogen exposure in the X70 would originate around the sulfide type inclusions because of their ability to irreversibly trap hydrogen.

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48 a b Figure 4.8 Ternary diagrams in wt pct of two inclusions families a Al-Ca-S and b Ca-Mn-S present in the X60 plate steel. Ternary diagrams were generated by automatic feature analysis at the Nucor Indiana facility. Dashed regions 1and 2 in a and b respectively show the grouping of the relative compositional distribution of inclusions evaluated by AFA in reference to the three elements found in each ternary diagram. Regions 3 in b indicate the presence of MnS type inclusions identified through AFA. a b Figure 4.9 Ternary diagrams in wt pct of two inclusions families a Al-Ca-S and b Ca-Mn-S present in the X70 plate steel. Ternary diagrams were generated by automatic feature analysis at the Nucor Indiana facility. Dashed regions 1and 2 in a and b respectively show the grouping of the relative compositional distribution of inclusions evaluated by AFA in reference to the three elements found in each ternary diagram. Regions 3 in b show the presence of MnS type inclusions identified through AFA.

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49 Figure 4.10 displays the ternary diagrams generated by AFA analysis. Within the two separate ternary plots a and b in Figure 4.10 the dashed areas regions 1 and 2 represent the common range of inclusion compositions. In region 1 the range inclusion compositions are Ca 25 – 75 wt pct S 35 – 75 wt pct and Al - 35 wt pct Figure 4.10 a. In region 2 the range of inclusion compositions are Ca 65 wt pct S 35 wt pct and Mn - 30 wt pct Figure 4.10b. The 100XF material had the highest amount of manganese 1.8 wt pct of all the materials but unlike the X60 and X70 at a similar level of sulfur 0.003 wt pct the inclusion compositions did not show a shift towards the manganese sulfur tie line. This is believed to be attributed to differences in steeling making operations that the 100XF underwent compared to the X60 and X70. Steel casting operations were not obtained from the steel producers. The 100XF exhibits low amounts of aluminum type nonmetallic inclusions in the steel matrix shown in region 3 of Figure 4.10a. It is speculated that if the inclusions within this area Figure 4.10a region 3 were evaluated with other elements i.e. Fe and O the composition of the inclusion would be different i.e. Fe-S and Al-O type inclusions. a b Figure 4.10 Ternary diagrams in wt pct of two inclusions families a Al-Ca-S and b Ca-Mn-S present in the 100XF plate steel. Ternary diagrams were generated by automatic feature analysis at the Nucor Indiana facility. Dashed regions 1and 2 in a and b respectively show the grouping of the relative compositional distribution of inclusions evaluated by AFA in reference to the three elements found in each ternary diagram. Regions 3 in a show a region where the types of inclusions within the area if evaluated with other elements i.e. Fe and O the composition of the inclusion would be found to be different. Speculated that these inclusions would be Fe-S and Al-O type inclusions. 4.2 Mechanical Properties Figures 4.11 – 4.14 show microhardness data from each rolled face to the center of each plate thickness along with the average microhardness solid line with 90 pct confidence interval calculated from the standard deviation in part a. Part b shows etched 2 pct Nital macrographs taken from the rolled face to middle thickness. Finally part c shows non-etched macrographs taken from the rolled face to mid-thickness. This analysis of

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50 hardness in part a was used to identify the presence and location of hard bands. The microstructure as a function of plate thickness was evaluated from the macrograph shown in part b. The non-etched specimens in part c reveal inclusion location and distribution. Figure 4.11shows the hardness traverse microstructure and inclusion distribution as a function of thickness for the X52steel. The hardness in the X52 decreases from each rolled face stabilizing to a constant value after approximately 8 pct through-thickness Figure 4.11a. There is also a slight rise in hardness from the average hardness 162 HV towards the centerline of material but otherwise there are no hard bands through the thickness.. The mostly consistent hardness through-thickness is consistent with the homogenous microstructure present from the edge to mid-thickness Figure 4.11b. The inclusions in X52 also appear to be distributed homogeneously Figure 4.11c. a b c Figure 4.11 a Microhardness data from each rolled face to the center of the X52 plate thickness along with the average microhardness solid line with 90 pct confidence interval b etched 2 pct Nital and c non-etched macrographs taken from the edge to middle thickness. Cross-sections of the full plate thickness of the transverse plane with respect to rolling direction were evaluated. Microhardness measurements were taken every 0.254 mm 0.01 in 500 gmf load and 10 s dwell time.

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51 Figures 4.12 and 4.13 show the hardness traverse microstructure and inclusion distribution as a function of thickness for the X60 and X70 alloys. The X60 and X70 show similarities with respect to the hardness traverse taken from each rolled face Figures 4.12a and 4.13a. Hardness is greater than average X60 – 202HV and X70 – 210HV from the rolled face to a thickness of 20 pct plate thickness after the quarter positions in each material the values are slightly lower than the average hardness. The X60 and X70 grade have very similar mechanical properties and microstructures that were produced during steelmaking and subsequent processing operations. There are higher amounts of banding and elongation towards the mid-thickness for the X70. The X60 and X70 exhibited no real observable hard bands based on the evaluation of Figures 4.12b and 4.13b or inhomogeneous inclusion distribution Figures 4.12c and 4.13c. Because a comprehensive study on inclusion distribution was not conducted during the thesis work areas of alloy segregation cannot be ruled out. a b c Figure 4.12 a Microhardness data from each edge to the center of the X60 plate thickness along with the average microhardness solid line with 90 pct confidence interval b etched 2 pct Nital and c non-etched macrographs taken from the edge to middle thickness. Cross-sections of the full plate thickness of the transverse plane with respect to rolling direction were evaluated. Microhardness measurements were taken every 0.254 mm 0.01 in 500 gmf load and 10 s dwell time.

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52 Figure 4.14 shows the hardness traverse microstructure and inclusion distribution as a function of thickness for the 100XF alloy. The 100XF edge 2 exhibited a higher hardness relative to the average hardness 280 HV for the first 20 pct through thickness of the plate Figure 4.14a. The higher values of hardness produce areas that would be more susceptible to HIC and thus the first 20 pct through thickness of the material might be more vulnerable to HIC damage. It was inferred the edge experienced higher cooling rates allowing for higher hardness to be achieved during rolling operations. The macrographs Figure 4.14b produced for the 100XF do not exhibit higher amounts of second phase microconstituents or drastic changes in microstructure due to plate thickness. Additionally inclusions appear to be equally distributed throughout the matrix Figure 4.14c. a b c Figure 4.13 a Microhardness data from each edge to the center of the X70 plate thickness along with the average microhardness solid line with 90 pct confidence interval b etched 2 pct Nital and c non-etched macrographs taken from the edge to middle thickness. Cross-sections of the full plate thickness of the transverse plane with respect to rolling direction were evaluated. Microhardness measurements were taken every 0.254 mm 0.01 in 500 gmf load and 10 s dwell time.

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53 a b c Figure 4.14 a Microhardness data from each edge to the center of the 100XF plate thickness along with the average microhardness solid line with 90 pct confidence interval b etched 2 pct Nital and c non-etched macrographs taken from the edge to middle thickness. Microhardness measurements were taken every 0.254 mm 0.01 in 500 gmf load and 10 s dwell time. Hardness measurements after prestrain were conducted as a measure of the amount of work hardening that occurred due to prestrain. Results from the hardness measurements are shown in Table 4.1. The hardness for the X52 increases the most due to the higher amounts of prestrain introduced increase from 157 to 200HV. At similar levels of prestrain 5 pct the X60 hardness increases more than the X70 increase of 22HV and 15HV respectively but the overall hardness at 5 pct prestrain is greater in the X70 218 HV than the X60 211 HV. The low prestrain amount for the 100XF does not increase the hardness. Hardness for plate steels used for sour service are is advised to be below 248 HV 3. Based on the hardness measurements HIC susceptibility is expected to be greatest in the 100XF condition.

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54 Table 4.1 Microhardness Data Taken for each Material and Prestrain Condition Material Prestrain Condition Microhard ness HV 90 Confidence HV X52 0 157 2 X52 12 200 4 X52 18 200 1 X60 0 190 1 X60 3 206 3 X60 5 211 3 X70 0 203 1 X70 5 218 3 X70 7 222 4 100XF 0 271 2 100XF 2 272 2 4.3 Characterization Study A characterization study on the fabricated double cell Figure 3.8 was conducted to optimize charging conditions. This study tested the ability of the test method to produce cracks in the material. The 100XF as-received plate was selected for the characterization study as it was assumed to have the highest HIC susceptibly. The high susceptibility allows for ease in identifying parameters that could alter charging conditions during the characterization study. Two variables 1 applied current density and 2 duration of test were varied to observe their effects on HIC damage the applied current density ranged from 0.80 to 25 mA/cm 2 duration set at 24 hours and charging times of 24 36 and 48 hours were assessed applied current density set at 15 mA/cm 2 . The experimental matrix is shown in Table 4.2. Current density was varied first while holding test duration constant. A total of two tests were conducted for each condition with the exception of the 25.00 mA/cm 2 condition due to deleterious conditions that occurred during charging which are described later. Once the optimal applied current density was identified it was held constant while test duration was varied. Three tests for each condition were conducted during this round of experiments to identify the influence of test duration. Table 4.2 Experimental Matrix for the Characterization Study Conducted on the 100XF alloy Applied Current Density mA/cm 2 Duration of Test hrs 0.80 1.50 5.00 10.00 15.00 25.00 24 2 2 2 2 3 1 36 - - - - 3 - 48 - - - - 3 - The numbers denote the number of tests conducted at each test duration and applied current density pairing After charging the specimens during applied current density experiments the specimen surface conditions were observed. A visual comparison of the surface conditions for the range of applied current densities is shown in Figure 4.15. The luster of the charged specimens is affected the applied current density. The most pronounced

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55 effects are observed in the 0.80 and 25 mA/cm 2 cases. The luster of these two specimens is far less than the other four. It is speculated that the 0.80 mA/cm 2 current density is associated with a lower cathodic-overpotential compared to the higher current density and as such galvanic corrosion of iron can occur at the metal/solution interface Figure 4.15a. A possible surface reaction at the highest current density 25 mA/cm 2 is also possible although the applied current density was sufficient to provide cathodic protection. At this current density the cell voltage and current changed significantly from the initial readings during testing. After 18 hours of EC charging the cell potential changed from 20.85 Volts to 28.51 Volts and the current dropped from 1.622 Amps to 0.725 Amps. Deterioration and mass loss of the steel sample were observed which subsequently polluted the electrolyte and is believed to be the cause of the change in potential Figure 4.15f. At intermediate current densities of 5 to 15 mA/cm 2 surface blisters appeared on the surface of the test specimens Figure 4.15c – e. The amount of blisters present on the surface of the 15 mA/cm 2 condition Figure 4.15e was greater than the other two applied current densities. a b c d e f Figure 4.15 Surface conditions of 100XF after cathodic charging at various applied current densities while keeping the duration of the test constant at 24 hours. a 0.80 b 1.50 c 5.0 d 10.0 e 15 and f 25 mA/cm 2 . Image taken using light optical flash photography. Upon completion of hydrogen charging specimen faces 2 4 and 6 Figure 3.9a were sectioned polished and evaluated for cracks. Face 2 is the furthest from the electrode connection as shown in Figure 3.6. The charging conditions that produced cracking for each of the faces and those that did not are shown in Table 4.3. For the lower current densities of 5.00 and 10 mA/cm 2 cracking was not observed throughout the entire specimen. An applied current density of 15 mA/cm 2 produced hydrogen cracks on each of the faces evaluated. The typical cracks observed are shown in Figure 4.16. The cracks generated were mainly long and straight but there were some regions of step- wise cracking. The CLR CTR and CSR are plotted as a function of applied current density in Figure 4.17. For the values of current density that were evaluated the crack ratios reached their maximum at 15 mA/cm 2 and then

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56 decreased at 25 mA/cm 2 . The values for the 25 mA/cm 2 current density condition are likely affected by the surface effects that were observed during testing Figure 4.15f. Table 4.4 shows the average crack depth measured from the top surface as a function of the applied current density. Most of the cracks occurred at depth of approximately 2 mm Figure 4.16 and very few samples exhibited cracks at the centerline 6.25 mm. The applied current density of 15 mA/cm 2 produced the largest range of crack depths. Because an applied current density of 15 mA/cm 2 produced the greatest extent of cracking it was chosen for the current density for further experiments. The effect of charging time on hydrogen cracking was explored by performing experiments at a range of charging-times while controlling current density 15 mA/cm 2 . The diffusivity of hydrogen in BCC ferrite is reported to be 10 -6 to 10 -4 cm 2 /s at room temperature 47 and 48. The hydrogen penetration for the different charging conditions can be estimated based on these diffusivity values. A simple estimate for diffusion distance lattice diffusion is 49: where x is the diffusion distance D z is the hydrogen diffusivity and t is the elapsed time in seconds during which diffusion occurred. The equation is an approximation since it is based on unidirectional diffusion in a system where the dimension in the diffusion direction is infinite. Cracks in the 100XF material were predominately within 3 mm of the outer exposed surface perhaps implying that damage deeper in the specimens was diffusion limited after charging for 24 hours. Using a diffusivity of 10 -6 cm 2 /s a lower limit the diffusion distance of hydrogen at room temperature is approximately 2.9 mm for 24 hours which approximately correlates to the value of crack depth observed at an applied current density of 15 mA/cm2 at the test duration of 24 hours Table 4.4. Table 4.5 shows crack depth data as charging time is changed. At longer charging times of 36 and 48 hours at 15 mA/cm 2 the depth of cracking did not change. Also electrolytic charging of the 100XF occasionally produced cracking close to the centerline 6.06 mm at 15 mA/cm 2 and 24 hours and thus it was concluded that hydrogen was able to diffuse through the thickness of the specimens within 24 hours. The crack depths observed for this material is believed to be related to the higher hardness in the first 20 pct through thickness of the plate Figure 4.14a. It is also believed that cracks that are generated in the first 24 hours inhibit the growth of other cracks deeper in the specimen. In support of this argument Griesche and Dabah et. al. 50 electrically charged duplex stainless steel samples and evaluated the presence of hydrogen though neutron imaging. The results from this study showed that sub-surface pores generated from the blister/void mechanism had high amounts of hydrogen compared to other areas found in the cross-section of the material. The pores were likely produced by the volume expansion produced through the formation of H 2 at grain boundaries producing internal stress that created the cavity. If hydrogen is trapped in the initially cracked regions HIC deeper in the specimen could be inhibited. The effect of test duration on the crack ratios is shown in Figure 4.18. The CLR at all test durations is greater than the CTR values. The scatter in the data for the critical crack ratios suggests that test duration does not significantly change the amount of cracking in the 100XF. √ 4.1

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57 Table 4.3 Charging conditions that produced cracking in 100XF that was cathodically charged for 24 hours Applied Current Density mA/cm 2 Face Examined 0.80 1.50 5.00 10.0 15.0 25.0 2 - - X X X X 4 - - - - X - 6 - - - - X - X Marks Presence of Cracks on Examined Cross-sections a b Figure 4.16 Hydrogen Induced Cracks produced through electrolytic charging of 100XF at a current density of 15 mA/cm2 for 24 hours aFace 2 b Face 4. Figure 4.17 Variation of calculated crack ratios versus applied current density for 100XF that was cathodically charged for 24 hours. CSR – Critical Size Ratio CLR – Critical Length Ratio and CTR – Critical Thickness Ratio.

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58 Table 4.4 Average crack depth in 100XF specimens at the current densities explored during EC experiments run for 24 hours: Extremes represent the minimum and maximum crack depth observed Average Thickness 12.5 mm Extremes mA/cm 2 Depth mm Standard Deviation mm Min mm Max mm 0.8 - - - - 1.5 - - - - 5 2.13 0.09 2.07 2.19 10 1.54 0.90 0.34 2.17 15 2.33 1.87 0.42 6.06 25 1.78 0.07 1.73 1.83 Table 4.5 Average crack depth in 100XF specimens as a function of test duration during EC experiments run at 15 mA/cm 2 : Extremes represent the minimum and maximum crack depth observed Average Thickness 12.5 mm Extremes Test Duration hrs. Depth mm Standard Deviation mm Min mm Max mm 24 2.33 1.87 0.42 6.06 36 2.12 0.87 0.64 2.64 48 2.17 0.68 0.68 2.60 Figure 4.18 Variation of average calculated crack ratios versus charging time for 100XF cathodically charged at an applied current density of 15 mA/cm 2 . 4.4 HIC Susceptibility Measurements This section presents the degree of cracking that was observed for each alloy and prestrain condition and discusses the results in light of differences in mechanical properties tensile strength and hardness.

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59 4.4.1 Influence of Prestrain HIC was observed in all but the X52 12 pct and X60 3 and 5 pct prestrain conditions exposed to the H 2 S method. Influence of Prestrain on Calculated Critical Crack Ratios Each pre-strained test specimen was charged with hydrogen either by the EC method or H 2 S method. The crack ratios for each of the two charging methods alloy and pre-strain condition are shown in Figure 4.19 - Figure 4.22. Error bars on the plots show the 90 pct confidence interval calculated based on the standard deviation of the population as well as the sample size. The results from the EC and H 2 S methods are comparable in magnitude to one another even though there is generally significant scatter in the data for both methods. The most significant difference is apparent for the X60 alloy Figure 4.20 which exhibited significant cracking when charged by the EC method but only minor cracking when charged by the H 2 S method. The differences in behavior are believed to be due to the statistical uncertainty of the evaluation method used. The statistical uncertainty of the evaluation method is possibly related to sectioning predetermined areas of the full specimen geometry. The sectioning process could miss crucial cracks that would contribute to the calculation of critical crack ratios. The critical crack ratios for the 100XF material and prestrain conditions are considerably greater compared to the other three alloys. a b Figure 4.19 Critical Crack Ratio values as a function of prestrain for the X52 material for each charging method: a Electrolytic charging and b H 2 S Method. CLR – Crack Length Ratio CTR – Crack Thickness Ratio and CSR – Crack Size Ratio.

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60 The HIC results are largely independent of pre-strain level within the statistical accuracy of the methods employed with the possible exceptions found in the X52 Figure 4.19a and X70 Figure 4.21a. The EC CLR results for X52 and X70 possibly exhibit an effect of pre-strain though HIC increases with increasing prestrain for the X70 while HIC decreases with pre-strain for the X52. It is speculated that dislocation density increases with increasing prestrain levels which might be expected to mitigate HIC by trapping hydrogen at dislocations 51 rather than more detrimental sites such as inclusions in the material. The X52 results are consistent with this hypothesis. However prestraining increases hardness and decreases ductility which might make steel more vulnerable to hydrogen embrittlement consistent with the X70 results. It is also worth noting that the prestrain levels of the X70 are lower than the pre-strain levels of the X52 and thus the increase in trapping sites is less in X70. However neither of the trends observed in the X52 and X70 alloys are observed in the other alloy conditions for either test method so the statistical certainty based on the evaluation method is unclear. The scatter observed with the critical crack ratios and prestrain is believed to be a reflection of the large variance associated with and intrinsic to the evaluation method employed. a b Figure 4.20 Critical Crack Ratio values as a function of prestrain for the X60 material for each charging method: a Electrolytic charging and b H 2 S Method. CLR – Crack Length Ratio CTR – Crack Thickness Ratio and CSR – Crack Size Ratio.

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61 a b Figure 4.21 Critical Crack Ratio values as a function of prestrain for the X70 material for each charging method: a Electrolytic charging and b H 2 S Method. CLR – Crack Length Ratio CTR – Crack Thickness Ratio and CSR – Crack Size Ratio. a b Figure 4.22 Critical Crack Ratio values as a function of prestrain for the 100XF material for each charging method: a Electrolytic charging and b H 2 S Method. CLR – Crack Length Ratio CTR – Crack Thickness Ratio and CSR – Crack Size Ratio.

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62 4.4.2 Influence of Mechanical Properties on HIC The HIC susceptibility of the non-strained specimens was compared based on the tensile strength. As the strength of the material increases the tolerance to hydrogen concentration generally decreases 14 and 15. Figure 4.23 shows the dependence of CLR and CTR values on tensile strength for non-strained test specimens exposed to both EC and H 2 S methodologies. Horizontal lines in Figure 4.23 represent the values for which steel plates are deemed fit for sour service in accordance with ISO 3183 52 this standard states that at values of CLR less than 0.15 Figure 4.23a and CTR less than 0.05 Figure 4.23b plate steel is qualified for sour service. The vertical line represents the NACE recommended maximum tensile strength 116 ksi to avoid hydrogen-assisted cracking phenomena. Figure 4.23a shows the relationship between the CLR values from each method as a function of tensile strength for the non-strained conditions. The CLR values presented in Figure 4.23a are consistent with the expectation that HIC increases with increased strength 14 and 15. The same trend is observed for the CTR values shown in Figure 4.23b. These results also show that the difference between the HIC resistance of the X52 X60 and X70 alloys from the 100XF alloy is approximately the same for samples charged by both the EC and H 2 S methods. Differentiation between the HIC resistance of the X52 X60 and X70 alloys is not evident with any statistical significance scale of 90 pct confidence intervals found in Figures 4.19 – 4.21 for results obtained from either method. Based on the results for both CLR and CTR the X52 X60 and X70 materials could be considered for sour service. a b Figure 4.23 a Crack length ratio dependence on tensile strength for the 0 pct prestrain condition: Electrolytic method EC and NACE Standard method H 2 S. b Crack thickness ratio dependence on tensile strength for the 0 pct prestrain condition. The vertical dashed lines represent the suggested 116 ksi threshold value for hydrogen assisted-cracking phenomena 47 and 48. The two horizontal dashed lines represent sour service requirements outlined by ISO 3183 52.

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63 Hardness measurements after the introduction of prestrain were performed to relate HIC susceptibility to the hardness change of each condition. Figure 4.24a and b shows the CLR and CTR values as a function of the measured hardness for each condition after prestraining. Neither crack ratio is strongly affected by hardness in the X52 X60 and X70 alloys. The ISO 3183 crack ratio values are plotted as horizontal dotted lines on Figure 4.24 and the vertical dotted lines show the suggested threshold hardness 22 HRC 248 HV where materials become susceptible to hydrogen embrittlement. As expected based on its higher strength and lower ductility the 100XF at 0 and 2 pct prestrain has crack ratio values that disqualify it for use in sour service 3 14 and 15. All of the lower strength alloys have lower hardness values than this threshold even after pre-straining which is consistent with their low HIC susceptibility. Their relatively low hardness may also explain why there is no apparent and consistent effect of pre-strain on their HIC susceptibility. a b Figure 4.24 a Crack length ratio and b crack thickness ratio dependence on hardness for the materials and pre-strain conditions evaluated by EC. The vertical dashed lines represent the suggested 22 HRC 248 HV threshold value for HIC susceptibility 47 and 48. The two horizontal dashed lines represent sour service requirements outlined by ISO 3183 52. 4.5 Hydrogen Analysis The following section presents the hydrogen contents found by both hydrogen analysis methodologies presented in section 3.7 LECO® and 3.8 mercury displacement. The results of total hydrogen concentration immediately after EC and values of trapped hydrogen in as-received and degassed specimens are presented in section 4.5.1. Values of diffusible hydrogen concentration directly calculated by the mercury displacement method and estimated using the LECO measurements and Equation 3.6 are presented in section 4.5.2. The LECO and mercury displacement method used different sample geometries and the influence of sample size on diffusible hydrogen results is discussed in section 4.5.3. HIC susceptibilities are then related to the diffusible hydrogen values and the amount of hydrogen permanently trapped in the material after exposure to EC.

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64 4.5.1 Average Hydrogen Contents Total and Trapped from LECO® Analysis Initial hydrogen concentrations in as-received plate steels were evaluated using the LECO® Hydrogen Analyzer the results are displayed in Figure 4.25a. The values represent the average of 5 measurements taken on each as-received plate steel. Error bars represent a 90 pct confidence interval calculated on the standard deviation of those 5 measurements. All of the alloys had levels of hydrogen less than 0.45 ppm Figure 4.25a. These levels suggest that the as-received alloys did not have hydrogen in high enough levels to produce HIC damage. Using the daisy chain setup described in section 0.1 total hydrogen concentration in all of the conditions was evaluated after exposure to the EC method the results are shown in Figure 4.25b. The values represent the average of 5 measurements taken for each material and prestrain condition. Error bars represent a 90 pct confidence interval calculated on the standard deviation of those 5 measurements. The total hydrogen concentration consists of both trapped and diffusible hydrogen. The total hydrogen present after EC varied between 2 and 6 ppm. The hydrogen concentration in the X52 and X70 increases with increasing prestrain level. The hydrogen concentrations of the X60 and 100XF alloys are independent of prestrain level. Hydrogen concentration is greatest for the 100XF 6 ppm and lowest for the X52 3 ppm. The comparison of these two materials represents the extremes of the materials investigated. These two materials in the as-received and prestrain conditions have drastically different levels of hardness 157 HV for X52 and 270 HV for 100XF and the microstructure was significantly different as well. Conversely the similarities of chemical chemistry and microstructure for the X60 and X70 allow for similar levels of hydrogen to be achieved 4 ppm. The effect of differences in microstructure on measured hydrogen contents is elaborated further in section 4.5.2. a b Figure 4.25 a Hydrogen concentration in as-received alloys. b Hydrogen concentration for each plate material as a function of prestrain.

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65 Diffusible hydrogen can move through the metal matrix and encounter microstructural constituents with high binding energies 60 kJ/mole and become irreversibly trapped 53. Values of binding energies for some typical hydrogen traps in steel are listed in Table 4.6. Depending on the type of trap HIC susceptibility has been shown to increase or decrease 12 and 54. Dislocations and grain boundaries are believed to exert a transient trapping effect on hydrogen as it moves through the lattice which was observed by Bouraoui et al. 55. However features such as inclusion and precipitate interfaces serve as irreversible trap sites. The average values obtained from the LECO® trapped hydrogen analysis for the materials and prestrain conditions are shown in Figure 4.26. The average values represent the average of 5 values. Error bars represent a 90 pct confidence interval that was calculated from the standard deviation of the 5 independent trapped hydrogen values. Other than X52 0 pct and 12 pct prestrain conditions values of trapped hydrogen were 1 ppm or greater Figure 4.26. The X52 trapped hydrogen concentration increases with prestrain. A statistically significant monotonic trend was not observed with respect to prestrain for the X60 X70 and 100XFalloys. Overall the trapped hydrogen in each condition is comparable indicating that there are no significant differences in irreversible trap sites between the alloys. The trapped hydrogen content of the alloys is likely mostly independent of pre-strain level because dislocations are reversible trap sites and thus increasing dislocation density through pre-straining should have minimal effect. Figure 4.26 Trapped hydrogen values after degassing of material. Materials were hydrogen charged with the EC method degassed at 50 °C for 72 hours sectioned and evaluated using the LECO® hydrogen analyzer to determine the trapped hydrogen content.

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66 Table 4.6 Hydrogen trap sites found in iron table modified from 54 Trapping site Binding/activation energy kJ/mol Assessment Method Reference H - dislocation 26 HTD analysis 56 H - grain boundary 49 HTD analysis 57 H - MnS 72 HTD analysis 58 H - Fe 3 C 84 Permeation 59 H - TiC interface 95 Permeation 60 4.5.2 Average Diffusible Hydrogen Content from LECO® and Mercury Displacement Analysis The diffusible hydrogen concentration was measured using the mercury displacement method on full size samples outlined in section 3.8.1 that were subjected to the EC method. Diffusible hydrogen was determined by averaging values from each of the four samples and a 90 pct confidence interval was calculated based on the standard deviation. The values for diffusible hydrogen determined from the mercury displacement method are shown in Figure 4.27a. The measured values of diffusible hydrogen for each material and prestrain condition showed high amounts of deviation. The greatest amount of diffusible hydrogen was observed in the 100XF material and prestrain condition 7 ppm and the least in the X52 material and prestrain conditions 2.5 ppm Figure 4.27a. The amount of deviation was attributed to non-uniform charging over the sample geometry which is discussed in the following section. Figure 4.27b shows the values of diffusible hydrogen calculated Equation 3.6 from the two values H Total and H Trapped obtained by the LECO Figure 4.25b and Figure 4.26. Error bars on the plot represent a 90 pct confidence interval calculated on the standard deviation of the 5 values. The levels of diffusible hydrogen determined by the LECO® method are depending on the material 1 to 3 ppm less than those measured with the mercury displacement method. This is likely due to the sub-sized specimens. The sub-sized specimens are believed to have more uniform charging during EC due to the smaller size represented by the tighter confidence interval Figure 4.27b. However the total hydrogen concentration values used in the calculation could be reduced during the lag time between placing each sample into the analyzer and subsequent melting and analysis by the LECO® analyzer 90 seconds. The time where the sub-sized specimens are at room temperature allows for a degree of hydrogen diffusion out of the sample. This reduction in measured total hydrogen level could also reduce the estimated diffusible hydrogen concentration. In both methods the diffusible hydrogen concentration is greater than 1 ppm for all conditions. Kittel et al. 61 suggested that a critical diffusible hydrogen concentration level of 1 ppm is necessary to induce damage within a steel matrix. These levels indicate that hydrogen was present in sufficient amounts to initiate HIC in all of the conditions Also the average values of diffusible hydrogen increase with increasing values of prestrain except for the X60 material. This trend agrees with literature 62 and 63 that at higher degrees of strains i.e. increases in dislocation density the diffusible hydrogen is increased. Bouraoui et al. 55 stated that dislocations exert a transient trapping effect on hydrogen atoms decreasing the hydrogen permeability though high purity iron and a similar effect is likely observed in the present study. The differences in the amount of diffusible hydrogen in each alloy could be related to the grain size and morphology present in the materials. The X52 is dominated by larger

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67 polygonal ferrite grains whereas the other alloys are dominated by finer acicular ferrite grains and probably a higher dislocation density. Of the acicular ferrite microstructures the amount of second phase microconstituents is highest in the 100XF alloy. It is hypothesized that the finer grain size larger grain boundary surface area higher dislocation density and higher surface area of boundaries at secondary microconstituents all serve to increase the transient trapping sites and result in the 100XF having the highest diffusible hydrogen content. a b Figure 4.27 Diffusible hydrogen values determined by two different evaluation methods on the four materials and prestrain conditions investigated: a Mercury displacement and b LECO® analysis. LECO® diffusible hydrogen values were determined using Equation 4.2. 4.5.3 Influence of Sample Location on Diffusible Hydrogen Results Sample Size Effect During the EC tests it was observed that gas bubbles nucleate grow and detach preferentially at the bottom end of the test specimen working electrode. This characteristic is displayed in the photograph shown in Figure 4.28. The evolution of bubbles was predominately within the first 25 mm of the specimen being charged sample 1 area shown in Figure 4.28. The effects of preferential bubble formation were assessed based on differences in diffusible hydrogen content as a function of sample number measured by mercury displacement and CLR value as a function of location of the examined face. Diffusible hydrogen measurements for samples 1 – 4 from each material and prestrain condition were compiled to evaluate possible non-uniform charging conditions. Figure 4.29 shows the dependence of diffusible hydrogen measured by the mercury displacement method with respect to the sample number referenced to the location of the electrode connection shown in Figure 4.28. The diffusible hydrogen amount in sample 1 for all materials is greater than samples 3 and 4. The amount of diffusible hydrogen decreases from sample 1 to sample

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68 4 which is evidence of the non-uniform charging. The non-uniform charging of these samples is believed to be due to a slight potential gradient on the sample during charging causing the evolution of hydrogen to be greater at the end of the sample. Because of the large difference in diffusible hydrogen content from sample 1 to 4 the standard deviation between the four values was also large producing the wide confidence interval in the diffusible hydrogen results shown in Figure 4.27a. The sub-sized specimens used for the LECO analysis are believed to have more uniform charging conditions which is indicated by the tighter confidence interval in the diffusible hydrogen estimates shown in Figure 4.27b. The CLR values from faces 1 2 4 and 6 of full-sized EC specimens were averaged together for each face to produce a single CLR value at the specific sectioning location this was done for each alloy and pre-strain condition. The CLR dependence on the examined face location is presented in Figure 4.30. Faces 1 and 2 are within areas of preferential bubble-formation and face 6 is farthest from it. The X70 Figure 4.30c and 100XF Figure 4.30d samples exhibit higher CLR values near the preferential hydrogen bubble formation region. The decrease in CLR value towards the top of the electrode i.e. from face 1 and 2 to the other faces face 4 and 6 can be interpreted as an indication that the presence of these bubbles at the bottom of the specimen corresponds to higher amounts of hydrogen damage in these regions. The measured crack ratios do not depend as greatly on location in the X52 Figure 4.30a and X60 Figure 4.30b samples which may be a reflection of the large variance in CLR associated with and intrinsic to the evaluation method employed. Figure 4.28 Observation of hydrogen bubbles forming preferentially on the lower bottom end of the steel specimens during EC experiments. Sample numbers are shown on the steel sample in the image.

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69 Figure 4.29 Dependence of location of sample sample number on the amount of diffusible hydrogen measured by the mercury displacement method. Sample 1 represents the sample furthest away from the electrode connection. Dimensions of the sample were the full thickness of the plate 100 XF and X70 - 12.7 X60 - 9.5 and X52 - 19 mm x 20 ± 3 mm width x 25 ± 2 mm length.

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70 a b c d Figure 4.30 Dependence of calculated Crack Length Ratio on the location of the examined faces for each material and prestrain condition for EC experiments a X52 b X60 c X70 and d 100XF. 4.5.4 HIC Susceptibility and Diffusible and Trapped Hydrogen Contents from LECO® Analysis Figure 4.31 shows the dependence of CLR as function of a the diffusible hydrogen content determined by the LECO® analysis method and b the amount of trapped hydrogen after degassing. HIC appears to be more dependent on the diffusible hydrogen concentration than the total trapped hydrogen concentration Figure 4.31a.

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71 Two key observations can be made from Figure 4.31a. The first observation is that for the highest strength material 100XF both CLR and diffusible hydrogen were the greatest. The CLR is believed to be associated with the higher strength and hardness of the material leading to higher amounts of cracking produced through the interaction with hydrogen whereas the higher diffusible hydrogen value is believed to be associated with increased surface area of grain boundaries and dislocation density to increase the overall transient trapping and therefore increasing the diffusible hydrogen. The second observation is that even with increases in the amount of diffusible hydrogen in the X60 and X70 acicular ferrite steels compared to the X52 alloy the degree of cracking in the 3 alloys is approximately equal. It would be expected from interpretation of literature 61 that at higher amounts of diffusible hydrogen HIC would increase. This trend is not observed and thus it can be inferred that for moderate and low strength steels changes in microstructure will not greatly impact HIC susceptibility. Figure 4.31b shows that HIC is independent of the trapped hydrogen amount in the conditions assessed. At similar levels of trapped hydrogen the HIC susceptibility of the 100XF is much greater than the other alloys. a b Figure 4.31 HIC CLR dependence on a LECO® diffusible and b trapped hydrogen contents from EC. The two horizontal dashed lines represent sour service requirements outlined by ISO 3183 CLR 0.15 52. 4.6 Evaluation of HIC with respect to Microstructure and Nonmetallic Inclusions Figure 4.32 shows representative images of the cracking behavior in each of the four alloys. Centerline cracking was observed for some of the sectioned faces for each material. Centerline cracking was observed more predominantly in the X52 Figure 4.32a X60 Figure 4.32b and X70 Figure 4.32c conditions whereas the 100XF material produced cracking near the rolled face average depth of 1.65 mm for all prestrain conditions Figure 4.32d consistent with the results obtained during the characterization study. The location of cracking within the 100XF alloy is believed to be due to higher levels of hardness compared to other areas within the plate thickness. The large range of crack depths shown in Table 4.7 observed for the X52 X60 and X70 implies that

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72 crack nucleation sites i.e. nonmetallic inclusions are distributed throughout the plate thickness as presented by the non-etched micrographs shown in Figures 4.11 – 4.13. The average crack depth shown in Table 4.7 represents the average of all cracks observed for the material and prestrain conditions evaluated. Each material had cracks at the centerline of the plate as well as cracks near the rolled surface. Table 4.7 Crack Depth Observed on all Sectioned Faces for each alloy after EC Material Half Thickness mm Average Depth of Cracking mm Standard Deviation mm Extremes Max Depth mm Min Depth mm X52 9.50 8.02 2.03 9.19 2.05 X60 4.75 3.02 1.27 4.54 0.59 X70 6.50 4.09 1.77 6.23 0.54 100XF 6.50 1.65 1.45 6.38 0.27 a b c d Figure 4.32 Light optical macrographs taken on as-polished full width and thickness EC test specimens. Images are representative of the cracking behavior observed for materials at all prestrain levels. a X52 0 pct prestrain Face 1 b X60 3 pct prestrain Face 2 c X70 5 pct prestrain Face 1 and d 100XF 0 pct prestrain Face 2.

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73 ESEM FESEM and EDS mapping were used to investigate the interactions of HIC with microstructural features such as parent microstructure secondary microconstituents/phases and nonmetallic inclusions. The results presented in the following section were obtained from 5 faces that showed the highest amount of cracking i.e. greatest CLR and CTR for each material. Primary cracking is denoted as the largest cracks whereas secondary cracking is the qualitative observation of cracking that branches from the largest cracks. Sectioned faces with the highest degree of cracking in X52 were evaluated after etching as shown in Figure 4.33. The cracking behavior had evidence of transgranular propagation though both the ferrite and pearlite Figure 4.33 features 1 and 3 as well as intergranular cracking of ferrite Figure 4.33 feature 2. Samples were back polished to 1 µm surface finish after being etched to conduct EDS on non-etched samples. Cracks on the polished surface are shown in Figure 4.34. Cracking was not observed around inclusions on the plane of view for the faces evaluated for HIC in the X52 this does not rule out the possibility of inclusions being present in the cracks out of the plane of view that was examined. If inclusions were observed they were outside the primary crack area Figure 4.34. The composition of the inclusions found within close proximity of the primary crack identified though EDS mapping consistently contained Al Ca and O. These are likely alumina inclusions encapsulated by a calcium sulfide shell as shown in Figure 2.5. Figure 4.33 Secondary electron micrograph taken on the FESEM of the X52 EC 18 prestrain Face 1 condition. Etched with 2 pct nital. Evidence of transgranular 1and 3 and intergranular crack propagation 2.

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74 Figure 4.34 EDS mapping of non-etched X52 EC 18 prestrain Face 1condition. Primary crack propagation in close proximity to Al-O-Ca type inclusion. SEM image produced was taken in backscatter mode on the FESEM see pdf version for color. The X60 and X70 steels were first etched with 2 pct nital to reveal the microstructure and then placed into the SEM to analyze the crack interaction with the microstructure. Figure 4.35 and Figure 4.36 show secondary electron images on the etched surfaces around the primary crack region for the X60 and X70 respectively. Images taken in secondary electron mode on the FESEM revealed that in the X60 Figure 4.35 and X70 Figure 4.36 HIC was associated with inclusions. It was difficult to determine whether the crack growth was transgranular or intergranular in nature but cracking associated with secondary microconstituents was occasionally observed. Figure 4.35 shows evidence of the primary crack propagating directly through two nonmetallic inclusions feature 1 and cracking that originates around the interface of an elongated nonmetallic inclusion feature 2 in the X60. Figure 4.36 also shows an example of HIC that propagates directly though a globular type nonmetallic inclusion in X70. After the microstructure was evaluated samples were polished to a 1 µm to specifically examine how inclusions interact with HIC. Using backscatter mode on the SEM gives composition contrast areas where nonmetallic inclusions were present around the primary crack region were identified. Once identified the areas were evaluated by EDS mapping to identify the type of inclusion present. Figures 4.37 and 4.38 show EDS mapping of primary crack regions where nonmetallic inclusions are present. In both SEM images there is a cluster of nonmetallic inclusions present in and around the primary crack area. EDS showed these clusters as being high in sulfur and manganese. MnS type inclusions were found by both EDS mapping and the AFA analysis presented in section 4.1. It is well documented in literature that these types of inclusions are detrimental to HIC resistance 3 5 8 and 18. This behavior is confirmed and observed for the X60 and X70.

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75 Figure 4.35 Secondary electron micrographs taken on the X60 EC 0 prestrain Face 6 condition. Image captured using the ESEM. Etched with 2 pct nital. Features 1 and 2 show HIC around inclusions in the microstructure. Figure 4.36 Secondary electron micrograph taken on the X70 EC 5 prestrain Face 1 condition. Image shows the presence of an inclusion in the primary crack area. Image was captured using the ESEM. Etched with 2 pct nital.

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76 Figure 4.37 EDS mapping of non-etched X60 EC 0 prestrain Face 6 condition. SEM image was taken in backscatter mode on the FESEM see pdf version for color. The SEM image and EDS maps show the presence of MnS type inclusions in and distributed around the primary crack region. Figure 4.38 EDS mapping of non-etched X70 EC 5 prestrain Face 1 condition. SEM image was taken in backscatter mode on the FESEM see pdf version for color. The SEM image and EDS maps show the presence of MnS type inclusions in and distributed around the primary crack region. Figure 4.39 shows secondary electron images from the 100XF alloy etched with 2 pct nital the images were taken using the FESEM. Cracking behavior in the 100XF showed: transgranular cracking of acicular ferrite grains Figure 4.39a feature 1 cracking of secondary microconstituents Figure 4.39b feature 2 and cracking around nonmetallic inclusions Figure 4.39c feature 3. At areas where nonmetallic inclusions were present the degree of secondary cracking decreases. In the absence of nonmetallic inclusions smaller micro-cracks branch off from the main crack increasing the degree of secondary cracking Figure 4.39a and b. It is speculated that the areas in the immediate vicinity of inclusions contains a high concentration of hydrogen so cracking is localized in these regions. Hydrogen is more uniformly distributed in regions without inclusions and therefore secondary cracking can occur more easily.

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77 Cracking behavior associated with inclusions in the 100XF is further shown in Figures 4.40 and 4.41. EDS mapping was conducted on both conditions to assess the type of inclusion present. Figure 4.40 exhibits a primary crack that propagates near but not through a spherical nonmetallic inclusion Al-O-Ca-S. Conversely in Figure 4.41 the primary crack interacts directly with the nonmetallic inclusion Mg-Al-O. It was observed that when the inclusions are more spherical in shape and are of the composition like the inclusion present in Figure 4.40 Al-O-Ca- S cracks do not interact with the inclusion. The slight elongation of the inclusion in Figure 4.41is likely more detrimental for HIC since there is a larger surface area in one plane for hydrogen accumulation and crack nucleation. The observed effects of inclusion shape as well as type of inclusion correlates with literature findings 3 18 21 and 31. Introducing CaS type inclusions or having a higher degree of spherical shaped nonmetallic inclusions from calcium treatment increases the overall HIC resistance with respect to inclusion susceptibility. a b c Figure 4.39 Secondary electron images take on a 100XF EC 0 prestrain Face 2 feature 1 shows transgranular cracking of acicular ferrite b 100XF EC 0 prestrain Face 2 feature 2 shows crack propagation across secondary microconstituent and c 100XF EC 2 prestrain Face 1 feature 3 shows nonmetallic inclusion. Image was captured using the FESEM. Etched with 2 pct nital. The interaction with microstructural features for each alloy is summarized in Table 4.8. The X60 and X70 interacted with hydrogen similarly. HIC appeared to be related to the presence of nonmetallic inclusions MnS in these alloys. The X52 and 100XF were the only alloys that showed signs of both intergranular/transgranular cracking within the fields of view observed as well as consistently observed cracking around secondary microconstituents pearlite in the X52 and martensite in the 100 XF. The X52 was the only material that did not show significant evidence of cracking around nonmetallic inclusions which is believed to be connected to the low sulfur content achieved during steelmaking lower manganese level and effective shape control over nonmetallic inclusions due to Ca additions. The X60 and X70 had higher amounts of manganese and sulfur 1.03 and 0.0007 wt pct respectively allowing for the formation of MnS and leading to preferential cracking around those features. Further characterization and analysis is needed on the X60 and X70 alloys to evaluate for the intergranular and transgranular cracking behavior. The 100XF had signs of multiple types of cracking: intergranular and transgranular cracking through ferrite grains cracking around secondary microconstituents martensite secondary cracking and cracking along nonmetallic inclusions oxides.

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78 Figure 4.40 EDS mapping of non-etched 100XF EC 0 prestrain Face 2 condition. SEM image was taken in backscatter mode on the FESEM see pdf version for color. SEM image reveals an globular shaped inclusion found outside the primary crack region. EDS mapping identifies the inclusion being of the mixed composition of Al-O-Ca-S. Figure 4.41 EDS mapping of etched 100XF EC 2 prestrain Face 1 condition. SEM image was taken in backscatter mode on the FESEM. Etched with 2 pct nital see pdf version for color. SEM image reveals an elongated inclusion found within the primary crack region. EDS mapping identifies the inclusion being of the mixed composition of Al-Mg-O. Table 4.8 Summarized Cracking Behavior Material Ferrite Secondary Microconstituents Secondary Cracking Nonmetallic Inclusions Intergranular Transgranular X52 X X X pearlite - -. X60 - - X - X sulfides X70 - - X - X sulfides 100XF X X X martensite X X oxides X - Presence of behavior “-“ denotes that further analysis is needed to determine presence of behavior

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79 CHAPTER 5: Summary and Conclusions Conclusions for the study of hydrogen induced damage in plate steels are discussed below in reference to the research objectives outlined in Chapter 1. 5.1 Electrolytic Charging A new dual cell electrolytic charging cell was designed for HIC experiments. The design addressed issues discovered in the previously used monocell design. The important features of the new electrolytic cell are the salt bridge to avoid cross contamination of the ion species from the cathode and anode a gas purge system to help agitate the solutions removal of oxide scale on steel specimens to inhibit oxide particles from contaminating the solution and an improved electrode connection. The new design was successfully used for stable long duration e.g. 24 hours experiments. A characterization study on the fabricated dual cell using the 100XF alloy studied the influences of applied current density and test duration to on HIC damage. It was found that at an applied current density 15 mA/cm 2 and a test duration 24 hours provided consistent HIC results. These testing parameters were then used for an investigation of HIC on non-strained and prestrained plate steels X52 X60 X70 and 100XF. The results were compared to NACE Standard TM0284 tests solution A on the same conditions. The critical crack ratios calculated from the EC method were similar in magnitude to the critical crack ratio values from the NACE test. Thus the EC method is possibly a viable substitute for the NACE standard test to investigate HIC. However further experimental refinement is necessary. The diffusible hydrogen values in the lower section of the specimen furthest section from electrode connection were significantly larger than the rest of the specimen there was approximately a 6 ppm difference in the 100XF and 3 ppm difference in the X52 alloy. The cause of the difference in hydrogen content is believed to be associated with non-uniform hydrogen charging over the specimen geometry. 5.2 HIC Susceptibility Measurements 5.2.1 Influence of Mechanical Properties The lower strength steels X52 X60 and X70 always have HIC susceptibility less than the ISO standard guidelines for qualification of sour service plates: values of CLR 0.15 and CTR 0.05. The 100XF regardless of prestrain level had the highest values of crack parameters: CLR 0.58 ± 0.06 CTR 0.07 ± 0.01 CSR 0.03 ± 0.005. The high susceptibility was attributed to higher tensile strength hardness and lower ductility leading to high amounts of HIC. These observations are consistent with NACE recommendations that HIC resistant alloys have tensile strength and hardness less than 116 ksi 800 MPa and 248 HV 22 HRC.

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80 5.2.2 Influence of Uniaxial Prestrain The HIC results are largely independent of the pre-strain levels imposed within the statistical accuracy of the evaluation methods employed. HIC resistance was largely independent of the effects of prestrain including increases in hardness dislocation density and reductions in ductility. Further prestrain did not increase the level of HIC in the X52 X60 and X70 alloys to a point where it would be disqualified for sour service according to the ISO standard guidelines. 5.3 Hydrogen Analysis The total amount of hydrogen which includes both the diffusible and trapped hydrogen was greatest in the 100XF alloy conditions and least in the X52 alloy conditions. The level of trapped hydrogen in the non-strained conditions did not vary greatly: X52 – 0.95 X60 1.07 X70 – 1.25 and 100XF - 1.18 ppm. This observation implies that the number of irreversible trap sites in the alloys is comparable. The differences in total hydrogen content are attributed to differences in the amount of diffusible hydrogen in the alloys. The average values for diffusible hydrogen are greatest in the 100XF 7 ppm and lowest in the X52 2.5 ppm values from mercury displacement. The large amount of diffusible hydrogen in the 100XF alloy compared to the X52 alloy is likely due to the relatively high grain boundary area and dislocation density. 5.4 Evaluation of HIC with respect to Microstructure and Nonmetallic Inclusions The qualitative analysis of microstructure and nonmetallic inclusions produced results that confirmed findings from literature. Cracking is observed around nonmetallic inclusions such as sulfides and oxides in the metal matrix. For materials in which both types are present X60 and X70 HIC is linked to areas of sulfide type inclusions. Elongated sulfide inclusions show the propensity to increase the degree of cracking.

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81 CHAPTER 6: Future Work This section outlines proposed areas of future work to further optimize the EC test apparatus and to further evaluate HIC in plate steels. It was observed that the overall size of the sample used in EC influenced the uniformity of charging that occurred during testing. It is proposed to test the effects of size/volume of EC specimens. The goal would be to optimize EC specimens for hydrogen charging uniformity. Because this study focused on the effects of uniaxial prestrain on HIC susceptibility a more comprehensive study involving different levels of residual stresses/strains to alter HIC behavior is suggested. This study should be applicable to actual industry practices used for plate forming operations that may alter plate steel resistance to HIC though the formation of residual stresses/strains. Conducting a study focusing on the microstructure influence on HIC susceptibility would be of interest. Targeted microstructures could be created by varying heat treatments cooling techniques and subsequent processing operations. The targeted strengths of the produced microstructures would have to be similar to be able to ignore the effects of increased strength to alter HIC susceptibility. The current evaluation method for HIC used in this study produced large variations in the cracking parameters calculated. The large variation made it difficult to discern differences between alloy and prestrain conditions. Similar concerns have been stated in the literature about the NACE standard H 2 S test. The sectioning procedure may miss cracks that could significantly change the test results. It is suggested that a study involving different evaluation methods be considered to increase the reliability of the results from HIC. The evaluation methods of interest could be Ultra-Sonic evaluation or 3-D X-Ray Tomography. These methods incorporate taking scans over the full size of the sample to identify defects such as internal cracks. It is thought that this type of evaluation may reduce the statistical error of the crack measurements. Inclusions in the low strength steels X60 and X70 provided the weakest link within the microstructure with respect to HIC. It was suggested by industry members that a more in depth inclusion analysis would be worth pursuing. This study would incorporate quantifying relative inclusion content found in the steel size distribution in the steel morphology produced and compositional mapping. Quantifying these variables would give rise to a rating scale of HIC susceptibility based on the features of the inclusion present within the steel.

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82 References Cited 1. J. Nieto T. Elias G. Lopez G. Campos F. Lopez R. Garcia and A.K. De “Development of Technology for the Production of HIC Resistant Slabs for Sour Service Applications at ArcelorMittal Lazaro Cardenas Mexico” Materials Science and Technology MST 2012. 2. T. Taira Y. Kobayashi H. Inagaki and T Watanabe “Sulfide Corrosion Cracking of Line-Plate or Sour Gas Service” NACE National Conference unpublished paper No. 171. Atlanta 1979. 3. G. Biefer “The Stepwise Cracking of Line-Plate Steels in Sour Envionments” In R. Kane R. Horvath and M. Cayard Wet H 2 S Cracking of Carbon Steels and Weldments pp. 479 - 494. NACE International 1996. 4. G. Herbslab R. Popperling and W. Schwenk “Occurrence and Prevention of Hydrogen-Induced Stepwise Cracking and Stress Corrosion Cracking of Low Alloy Pipeline Steels” NACE National Conference unpublished paper No. 9. Chicago 1980. 5. Y. Shinohara and T. Hara “Metallurgical Design of UOE Line Plate for Sour Service” Microalloyed Steels for Sour Service International Seminar São Paulo Brazil August 2012. 6. J. A. Ronevich “Hydrogen Embrittlement in Advanced High Strength Steels” M.S. Thesis Metallurgical and Materials Engineering Colorado School of Mines Golden CO 2009. 7. B. Rosner “Exploration of a Hydrogen Charging Method to Understand Performance of Energy Tubular Products” M.S. Thesis Metallurgical and Materials Engineering Colorado School of Mines Golden CO 2011. 8. D. Jones Principals And Prevention Of Corrosion 2nd ed. New Jersey: Prentice-Hall Inc pp. 5-18 40-50 75-96 143-148 249-254 279-286 335-342 439-444. 1996. 9. R. Kane S. Wilhelm and J. Oldfield “Review of Hydrogen Induced Cracking of Steels in Wet H2S Refinery Service” In R. Kane R. Horvath and M. Cayard Wet H 2 S Cracking of Carbon Steels and Weldments pp. 173 - 185. NACE International 1996. 10. M. Cayard and R. Kane “Characterization and Monitoring of Cracking of Steel Equipment in Wet H 2 S Service” In R. Kane R. Horvath and M. Cayard Wet H 2 S Cracking of Carbon Steels and Weldments pp. 633 - 651. NACE International 1996. 11. G. Krauss Steels: Processing Structure And Performance 4th ed. Ohio: ASM International pp. 101-116. 2005. 12. E. Moore and J. Warga Materials Performance vol. 15 No. 6 pp. 176. 1976. 13. A. Ikeda T. Kaneko and F. Terasaki “Influence of Environmental Conditions and Metallurgical Factors on Hydrogen Induced Cracking of Line Plate Steel” NACE National Conference unpublished paper No. 7. Chicago 1980. 14. S.U. Koh J.S. Kim B.Y. Yang and K.Y. Kim. “Effect of Line Plate Steel Microstructure on Susceptibility to Sulfide Stress Cracking” Corrosion Science vol. 60 No. 3. 2004.

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